Crystal structure and morphology of the carbide precipitated from martensitic high carbon steel during the first stage of tempering

CRYSTAL STRUCTU$IE AND MORPHOLOGY OF THE CARBIDE P~~~~P~TATE~ FROM MART~NS~T~~ HIGH CARBON STEEL DURING THE FIRST STAGE OF TEMPERINGIY.

HIROTSU$

and

S. NAQAKIJRAS

A ms,rtensitic steel containing 1.13 wt. %C was tempered at 120°C for l-100 days and the structure and morphology of precipitated carbide were studied by electron microscopy and selected area diffraction. The carbide is orthorhombic with lattice parameters a = 4.704 f 0.016, b = 4.318 If: 0.005 and c = 2.830 f 0.006 A. The space group is Pzrt*n and the atom positions tnw4Fe at 4g with 16 = 5, y = &and 2C at 2a. The carbide, naed ?I-crcr‘bide of iron or q-Fe&, is isomorphous with Co& and Co,N. The structure is simi~r to but different from the structure of hexagonal s-carbide, which has long been believed to precipitate at the fir& stageof tempering. The orientation relationshi between the q-carbide and the tempered marten&e EL”w&s determined as (llO)q 11(OR&* and [OOl]q ip[l@#b,-. The carbide has s pl&e-like shape with thickness of 30-50 A randprecipitates periodically along dislocations with an interval of about 100 A. STRUCTURE CRISTALLINE ET MORPHOLOGIE DU CARBURE PRECIPITE A PARTIR DE L’ACIER MARTENSITIQUE A HAUTE TENEUR EN CARBONE AU COURS DU PREMIER STADE DE RECTJIT Un &er martensitique contexmnt 1,13 % C en poids a 6th recuit 8, 120°C pour u.ned&e ds l-100 jours. Ls structure et le morphologie du oarbure p&&pit6 a Bte Btudi&sper microscopic Blectronique et per diffraction BUPune surface d&err&&e. Les auteurs trouvent que le carbum est orthorombique, bs paMImetres du r&au &ant: 4 = 4,784 f 0,016, 6 = 4,313 & 0,005 et c = 2,830 f 0,006 A* Le groupe d’aspace est Pnnm et 1%~positions des atomes sont 4Fe B 4g svec 1: = f, y = f, et 2C 8 2a. Le carbum, appel8 carbure q de for ou q-Fe&, eat isomorphe de Co& et Co,N. 58 structure est semblable, quoique diffkrente, de Ia structure du carbure E hexagonal, que l’on a longtemps suppose pr6cipiter au premier stie de recuit. La relation d’orientation entre le carbure 7 et le marten&a mcuite a” a 6th Le carbure se pr&sentesous forme de plad&ermin& comme &ant (110)~ 11(OlO)z- et [OOl]t) )I[lOOk. guettes de 30 $ 60 A d’epeisseur et precipite p&iodiquement le long des dislocations, avec des intervalles de 100 A environ, KRISTALLSTRUKTUR W&REND DER

UND MORPHOLOGIE DES AUSGESGHIEDENEN K&&BIDES ERSTEN ANLAOBEHANDLUNU VON MARTENSITI%HEM, K~~~NSTOFFREI~H~M STAHL Ein 1,13 Gew. % C enthaltender maicensitischer Stahl wurde bei 120% zwischen l-100 Tagen getempert und die Struktur und Morphofogie der Karbidaussoheidnngen mit Hilfe der Du.rchstmhlungsElektronenmikroskopie und der Feinbereichsbeugung untersucht. Das Karbid iet orthorhombisch mit den #ittepmmetern u = 4,704 -& 0,016, b = 4,318 f 0,005 und G = 2,830 & 0,006 A. Die Struktur gehort zur %umgruppe Pnnm und die Atomlagen sind 4Fe bei 4g mit o = & y = + und 2C bei 2a. Des Karbid, das TKrtrbid oder -Fe,C genennt wird, ist isomorph mit Co& und Co,N. Die Struktur ist 5hnlich wie die Struktur des% exagonalen e-Kerbids, von dem man hmge gegbubt hatte, d&l es sich im Anfangsstidium des Temperprozasses ausscheidet. Die Orientierungslmziehungzwischen tj-Karbid und dem getsmperten Martensit a” wurde bestimmt: (110)~ 11(OlO)c* und [OOlj, I/ [lOOk. Die Karbidausscheidungen sind plattenf”drmig mit Dicken zwischen 30 und 50 A und entlmg Versetzungen in etwa 100 A-Intervahen period&h angeordnet. f.

INTRODUCTION

The tempering of ferrous martensite has been the subject of many studies, and much information about the structural changes has been obtained from Kurdjumov and X-ray and electron diffraction. Lyssak(li firet pointed out, from X-ray studies, that during the first stage of tempering of ma&&tic high-carbon steels the psmary rna~nsi~ (a’-phase) transforms to a low carbon martensite (a”-phase) with an axial ratio of 1.012-l .013 and simultaneously a transitional carbide precipitates. This w&sconfirmed by Roberts et ~1.‘~) and a theory of the first stage of tempering was developed by Cohen et 02.(2-4) Concerning the transitional carbide coexisting with tempered martenaitn, Heidenreich eb &.‘b) reported, on the basis of an electron diffraction study, that the precipitates were the hexagonal nitside FesN or a t Received June 18, 1971. 8 Department of Metallurgy, Tokyo Institute of Technology, Oh-okayama, Meguro-ku, Tokyo, Japan. ACTA METALLURGICA,

VOL. 20, APRIL

1972

846

carbide isomorphous with it. Jack’“) made a more detailed study. He tempered martensitic steel of 1.3 wt. %c at 120°C for 5-40 days then examined it by an X-ray powder method. A carbide with h.c,p. structure was found and called e-carbide. This was the same as the synthesized h.c.p. carbide Pcs_sC by Hofer et al.f7) and by one of the present authors.@’ Jack also suggested an orientation rek+,tion&i~ between the s-carbide and the matrix, Using an X-ray single crystal technique, Arbuzov and Khayenko(s) derived the ofi&ation r&tionship similar to Jack’s. As is well known, selected area electron diffraction can be more advantageous than X-ray diffraction in studying fine p~cipi~~s in matrix phases. Thus, Tekjn and Kelly, fro) Well@) and recently Barton(rs~ studied the structure and the morphology of tmnsitional carbide in tempered Fe-Ni-C alloys, All of Y them found fine s-&bide particles lying across transformation twins and in untwinned regions.

ACTA

646

METALLURGICA,

Wells mentioned that the &-carbide, taking the form of thin plates or lath with a habit plane close to (100) of the matrix, precipitates frequently on more than one plane in any martensite plate. Tekin and Kelly observed the s-carbide precipitated coherently along (100) matrix directions. Barton showed dark field images of c-carbide in the shape of rods. In spite of its importance, no electron microscopy and diffraction seems to have been carried out for iron-carbon madnsitic steels subjected to lowtemperature tempering. Therefore, detailed studies have been made to clarify the crystal structure and the morphology of the transitional carbide. In contrast to the results of the previous works, it was found that the precipitated carbide was orthorhombic and isomorphous with CoZN and Co&. It precipitates coherently in the matrix with a fine plate-like shape. The lattice relationship between the carbide and the matrix was also determined. Some of the results have been reported previously. In this paper, a detailed description of the results is given. 2.

EXPERIMENTAL

PROCEDURES

A rod of zone-refined pure iron (Johnson-Mathey Chemicals Co. Ltd.) was cold rolled to a thickness of 0.13 mm and carburized at 950°C for 1 hr in a stream of town gas activated by passing through a converter furnace filled with charcoal. Then the specimens were austenitized at 1100°C for 7 hr in a high vacuum and quenched into ice water, followed by immersion in liquid nitrogen. A martensitic high carbon steel was obtained. The chemical analysis showed that the carbon content was 1.13 wt. ‘A. The concentrations of other elements are as follows: N: 0.006, Si:O.O02 and Mg, Mn, Cu, Al: <0.0003 wt. %. The specimens were tempered in an oil bath at 120°C for 1, 5, 10, 20 and 100 days. X-ray powder patterns of the specimens tempered for 10 and 100 days showed weak diffraction lines due to the precipitated carbide as well as lines due to the tempered martensite and retained au&mite. All spacings and intensities observed coincided with those given by Jack within experimental error. The tempered specimens were electrolytically thinned for electron microscopic observation, with H,PO,-CrO, electrolyte kept below 70°C. The electron microscope used is a JEM200 operated at 200 kV with a tilting stage.

VOL.

20,

1972

The twin plane is (112) .(l5) In agreement with this, a twinned structure was observed in almost all the martensite grains in the untempered specimen and it persisted even after the tempering for 100 days. Both in the twinned and untwinned regions, dislocations were observed with a high density of (l-4) x lOlo cm-2. Although their arrangement and density were not affected by the tempering, their image contrast was changed because of carbide precipitation along the dislocations (Section 3.4). The twinned structure makes the analysis of electron diffraction patterns complicated, because the diffraction patt’erns are usually composed not only of spots from the matrix and twins, but also of spots caused by the double diffraction between them. Figure 1 is a stereographic projection of the twinned structure. The solid circles correspond to the matrix lattice and the plain circles to the twin lattice. The double circles represent the orientations common to the matrix and the twin. The twin plane is chosen as (112), and (n2),, while the pole is (110),, and (liO),. Here, M and T indicate the matrix and the twin, respectively. Figure 1 is drawn for the tempered martensite with an axial ratio (c/a),- = 1.014, according to the experimental result described later. From the figure, it can be seen that specimens with the surface orientations represented by the coincident and nearly coincident circles give electron diffraction patterns consisting of two groups of diffraction spots,

3. RESULTS

3.1 Structure of tempered martensite High carbon steels subjected to the martensitic transformation are known to contain many fme transformation twins in plate-like martensite grains.cl4)

0 :

MATRIX(d”)

.‘I:

TWIN

Fro. 1. Stereographic projections of matrix and twin of tempered martensite with an axial ratio of 1.014.

IIIROTSU

AND

NAGAKTJRA:

CARBIDE

symmetrically or almost symmetrically arranged about B line passing through the indident spot. Such patterns will be called twinned diffraction patterns.(L*) In this study, it was found that the twinned ~ffr~ction patterns always contained the spots caused by double d.iffr&ion. When the specimen o~en~tion falls on the horizontal great circle in Fig. 1, the reilection from the twin plane (112), is always excited. Such patterns are useful in determining the unique axis c,~ and also the lattice relationship between the matrix and the precipitated carbide. An example is shown in Fig. 7(b). In this pattern, the diffrlaction spots from the matrix are accompany by streaks running along [f121X* directiont This is because the mertensite, even after tempering, contGns the repeated twin structure along the direction perpendicular to the twin plane. The repeated twin structure results in spikes from the reciprocal lattice points and the spikes sometimes make diffraction spots other than those expected from the specimen orientation appear in the ~~r&ction pattern. For example, as cttn be seen in Fig. 2(b), where the incident beam dire&ion is [lOOIM, the diffraction spots, T, corresponding to the reciprocal lattice points on (lTl)r* passing through the origin exist in addition to the spots from ( 100)M* plane, in spite of the angle between and (lIl),* being as Isrge ss 15’. The spot ( wH* due to the double diffraction indicated by arrow is also observed in this pattern, Accordingly, great care was taken to distinguish the spots due to the precipitated carbide from the spots due to the tempered martensite in anslysing the diffraction patterns, In order to determine the lattice parameters of the tempered martensite, TlCl was evaporated on the specimen surface and used a8 a reference. a”-200 and a”-002 reflections were excited as strongly 8s possible and their spacings were measured. Assuming the lattice parameter of TlCl to be 3.334 A, we obtained for the lattice parameters of the specimen tempered for one day a,* = 2.850 f

0.002 with

and

car = 2.891 f

(c/a&- = 1.014 f

PRECIPITATED

DURING

TEMPERING

647

tempered specimens. No essential difference was found emong the specimens subjected to various periods of tempering. Figures 2-9 reproduce the diffraction patterns with the incident beam direction being parallel to the im~rtant zone 8x88 of the matrix. The diffraction spots from the carbide (tre arranged regularly with respect to the matrix spots,t indicating that the carbide precipitates coherently within the matrix lattice. In these patterns, the spots indicated by the single arrows &re due to double diffraction between the matrix and the twin, while the spots indicated by

(a)

(b’

Fm. 2. Electron diffraction patterns: (a) beam j ETlO]*,]I [OOl],, tempered for 1 day; (b) beam14 1 I [OOl], 11 [lOO]x, tempered for 100 days. .iW, T and;q, matrix, twin and carbide, respectively. -c, double diffraction between matrix and twin. S, superlattice reflection.

-

Cal

_.T

fbl

FIU. 3. Electron diffraction patterns: (a) beam [Oll]~, tempered for 100 [Toll,, tempered for 1 day. matrix and carbide.

0.002 A 0.091

(1)

They agree well with those given by Kurdjumov axid Lyssak@) and by Roberts et ol.r2) The a,. value was almost independent of the period of tempering, but the c,. value showed a tindency to decrease slightly with increasing period of tempering. 3.2 Cr@ul structure of the precipitated carbide of

Nany single crystal electron diffraction patterns the p~cipi~~d carbide were obtsined from

f [hkZ]* and (wow)* represent the direction end the plane in the reciprocal lattice, respectively.

(al

(bl

Fro. 4. Electron diffraotion patterns: (a) beam [lOl]q [llT]x, tempered for 100 daye; (b) beam ii [Oil], II [Ill]~, tempered for 100 days. $ In what folfows, the word “matrix” is used without distinction between the matrix and the twin lattices.

ACTA

648

la)

METALLURGICA.

(b)

FIG. 5. Electron diffraution patterns: (a) beam tempered for 20 days; (b) beati [0211u, tempered for 1 day.

w&r9

VOL. 20, 1972

1

FXU.8. Electron diffraction pattern: [Zll]~,

(: a)

beam tempered for 100 days.

\I[OlZj, I\

CbJ

Fro. 6. Electron diffrsotion patterns: (a) becun j! [ITQ /j [Zol]~, tempered for 20 days; (b) beam 11[ITI], 11[lo&, tempered for 1 day.

(al

(b 1

FIG. 7. Electron difffsotion patterna: (a) beam j/ [lZl]q 11 [Tl33~, tempered for 1 day; (b) beam 11[OT3]q jj @ll&. tempered for 20 days.

the double arrows are due to double diffraction between the matrix and the precipitated carbide. It is noticed here that the incident beam directions bNi$f and [wvu],~ give different carbide patterns. Using these diffraction patterns, the reciprocal lattice of the precipice carbide was constructed in association with the reciprocal lattice of the matrix. In order to determine the correct positions of the reciprocal lattice points, the Bragg reflections from the precision were excited a~ strongly aa possible by the u8e of the tilting stage, and the effect of spikes from the reciprocal lattice points was eliminated.

Ir

Fra.9.Electron diffraction pattern:

beem [5311x, tempered for 20 dsye.

/1[126],, I/

The carbide net patterns can not be interpreted on the basis of the E-carbide with h.c.p. structure. According to the careful measurements of the lattice spacings, the diffraction patterns differed from the patterns expected from a crystal with sixth-fold symmet~. In addition to this fact, ~ff~etion qots unobtainable from the crystal structure of s-carbide appeared. In Figs. 2(a), 3(b), 5(a) and 6(b), such spots are indicated by the letter 8. Figure 10 represents a perspective view of the co~t~e~ reciprocal lattice. Here, the plain circles represent the reciproual lattice points for the carbide and the solid circles

HIROTSU

AKD

NAGAKURA:

CARBIDE

FXU. 10. Perspective view of the reciprocal lattices for the precipitated carbide and the matrix: open circles, carbide; filled circles, matrix; broken line, primitive reciprocal cell. S corresponds to superlattice reflection.

those for the matrix 01”. This result shows that the carbide is not hexagonal but orthorhombic. The broken lines indicates the primitive orthorhombic unit cell in reciprocal space. The lattice parameters of the carbide were determined by taking the lattice spacings of the matrix as standard, They are a,, = 4.704 & 0.016, and

b, = 4.318 & 0.005 c,, = 2.830 -& 0.006 A

(2)

We call, hereafter, this orthorhombic carbide qcarbide, for convenience. A comparison of the present unit cell with the unit cell of

s-carbide shows that a, N 2/!&z,, b, - c., and V,,N~V,, where a,(== 2.752 8) and c,(= 4.353 A) and V, are the lattice parameters and the unit cell volume of e-carbide, respectively.d7) This shows that the precipitated carbide contains four iron atoms in the unit cell. The q-carbide is pseudo-hexagonal. This is clearly shown in Fig,

cti-ae

PRECIPITATED

DURING

TEMPERING

3(a), in which the net pattern due to the carbide does not have hexagonal symmetry. Among the structures of transition metal carbides and nitrides, there are two types possessing such a pseudo-hexagonal unit cell. They are the Co,N type (CosN and Co&) and the <-Fe,N type. (r*) In these structures, the nonmetal atoms occupy regularly half of the octahedral interstices formed by the metal atoms, making a pseudo-hexagonal closed packed arrangement, although the occupied positions are different in the respective structures. As a result of the regular array of nonmetal atoms, superlattice reflections appear in diffraction patterns and their intensities are solely contributed by the nonmetal atoms. Corresponding to the difference in the nonmetal atom ar~ngements in the CoaN and <-Fe,N structures, superlattice spots appear in different positions in the diffraction patterns. In the present carbide, the spots indicated by the letter 8 in Figs. 2(a), 3(b), B(a) and 6(b) can be interpreted as the superlattioe spots. Figures 11(a)-(c) are other examples of diffraction patterns showing superlattice spots. The superlattice spot positions can only be in~rp~ted by assuming the CosN type of structure. Therefore, the q-carbide may have a structure isomorphous with Co,N and Co&, and may be Fe,C. The space group of q-carbide is Pnnm and the atom positions are 4 Fe at 4g with x = 3, y = i and 2 C at 2a. Figure 12(a) depiots the structure, with the solid and plain circles indicating carbon and iron atoms, respectively. In this structure, the carbon atoms are in the octahedral interstices. Figure 12(b) represents b,, projection of the structure. It is obvious that the iron atoms take a distorted h.c.p. arrangement and the int,erstitial carbon atoms line up on the straight

..I__

(a)

(bl

649

(cl

Cd)

Fm. 11. Electron diffraction patterns showing super-lattice reflections from the q-carbide: (a) tempered for 1 day; (b) 20 days; (c) 100 days. 8, superlattice reflection; F, fundamental reflection; dL[,T and q, matrix, twin and carbide, respectively; 4, double diffraction between matrix and twin. Patterns (a), fb) and (c) are obtained from almost the ssme beam incidence, beam j/ [ lli]~ jj [ 1201X, and their key diagram is drawn in (d).

ACTA

660

20,

VOL.

METALLURGICA,

1972

0

0

0

0

0

0

(b)

(a)

Fro. 12. (a) Structure of orthorhombic q-carbide. Atom positions are C: 0, 0, 0; 4, 4, it. Fe: !, 4, 4, 0; :9 4, _h, (b) btt projection of the structure: 0, carbon atom et 0; 0, carbon atom at hb; at fb; 0. iron atom at ab. lines

running

along

the

cB-axis with

an

TABLE 1. Spacings and intensi+.ieeof q-Fe&

interval of

a@ on b,-plane and an interval of b, on a,,-plane. In Table 1, calculated lattice spacings (da,) and intensities ([ Fcarlg, Pm, : calculated structure fact or) for q-Fe& are compared with observed spacings (dabs) and visual intensities (Iobs), respectively. The observed values coincide well with the calculated.

From Fig. 10, we oan determine the orientation relationship between the carbide and the matrix BS follows :

wQJ IIP1%” and PW, IIWI,.

(3)

This relationship can explain all the observed spot positions in the diffraction patterns. The indices of the ~ffraction spots and the reciprocal lattice points are given hereafter on the basis of this orient&ion relationship. Figure 13 is a stereographic projection showing the orientation relationship. The matrix is tetragonal tempered martensite with axial ratio (c/a),- = 1.014. The solid circles correspond to the matrix and the plain circle8 to the carbide. Figure 14 is a stereographic projection showing the orientation ~l&tionship between e-carbide snd a-iron proposed by Jack,(e) which is (lO.l), I( (lol),, (OO.l), 11(Oil), and 5’ from r.11.01,IIr1001,. Pit8ch and Schrader’e orientation relationship, whichis [ll.O], 11[NO],, [Tl.l], II[OlO], and [lI.l], II[OOI],, is similar to Jack’s, Comparing Jack’s and Pitsch and Schrsder’s orientation relationship with our own, we find a slight difference in

Claaeifieation* 8

Spanting (A)

Intensity

hkZ

doa

dabs

jFc’c.lf’t

110 101 011

3.18 2.42 2.37 2.35 2.16 2.11 2.06 1.96 1.67 1.61 1.69 1.47 1.42 1.39 1.38 1.37 1.31 1.29 1.28 1.27 1.24 1.23 1.21 1.19 1.17 1.17 1.16 1.15 1.13 1.13 1.08 1.06 1.06 1.05 1.03 1.02 0.86

3.17 2.42 2.37 2.36 2.16 2.12 2.06 I.;8 1.61 1.59

22.1 33.2 14.4 31.8 190.4 2i6.8 201.7 0 6.3 81.0 77.4

va VW 8 8

1.P2

18:::

;

i.6 169.6 0 3.1 3.0 0 24.0 23.6 10.9 64.3 66.3 10.6 60.2 0 60.7 1.8 136.0 28.7 1.4 1.4 27.0 1.3 49.6

1 VB uw m w VW 1: B w w -

:z 111 210 120 211 121 220 310 002 221 130 301 311 112 031 320 131 230 202 022 212 400 321 122 410 231 040 222 330 411 420 312 042

1.38 1.30 1.23 1.23 1.20 1.19 1.16 1.17 1.16 1.07 1.06 1.03 0.85

I obe WNVW .J 8” v.9 938

8

w _” ?n

* F and S refer to fundamental and super-lattice fleotioru, raspectivsly. t Relativistic correction is made. $ bfaalwd by matrix reflection.

re-

HIROTSU

0 0

0

&ND

q-urbm projected

NAGAKURA.:

e frwn

CARBIDE

Nt?ix oppositapolar

I%#. 13. Stereographic projection showing the orientation relationship between q-carbide and the matrix with an axial ratio (c&’ of 1.014.

PRECIPITATED

DURING

651

TEMPERING

pattern, fo) the dark field image and (d) the key diagram for (bf. The dark field image was taken with the q-101 spot. At first sight, dark feather-like images in (a) look like the precipitated carbides, but the dark field image (c) shows that the precipitated carbide particles are very small, about 50 x 100 A2 and line up fairly regularly along the dark feather-like images in (a). By tilting the specimen, the bright and dark field images were carefuliy compared. It was found that the feather-like images are due to dislocations, along which the carbide particles precipitated. When the carbide reflections are excited, dark feather-like or linear images appear usually in electron micrographs. Figures lS{a) and (b) are examples. In (a), where alternate white and black bands are due to the twinned structure, tine dark linear images are observed to run across the twin plates. Their direction is nearly parallel to the [ltl],8 direction projected onto the specimen surface, so they are considered to be twinning dislocations.(ls) In (b), where the twinned structure is not observed, ~sl~ationa or c&3-crossox) pattern. make a cross-hatched (a>

Fxo. 15. (a) Bright field image of a specimen tempered for 100 days, (b) corres onding diffraction pattern, (cf dark field image using q- P01 spot and (d) key diagram for the di&%actionpattern (b).

0 FKZ.

&-carbide

l

matrix(a-ironf

14. Stereographic prajection showing Jsck’s orientation relationship.

the stereographic proj~c~ons. This is due to the difference in the cry&al symmetries of c-Fe&2 and q-Fe&. 3.4 ~0~~~~~

of &iz ~~~~i~~~

abbe

In Fig. 15, (a) is a bright field image of the specimen tempered for 100 days, (b) the corresponding diffraction

0.5

(a)

w

0.5 CI

(b)

Fro. 16. Image of dislocations decorated by the preeipitated carbide particles. In (a), dark bands are due to the twins, In (b), dislocations form a criss-cross net.

ACTA

652

METALLURGICA,

All the dislocations are decorated by the precipitated carbide particles. Frequently, the dark linear images and cross-hatched patterns have been ascribed to images of carbide particles themselves, but this is not correct. Various dark-field images using carbide spots showed that q-carbide particles precipitate predominantly along dislocations with an interval of about 100 A. This is schematically shown in Fig. 17. Here, several carbide plates with the same shape are arranged in parallel with an interval D along the dislocation line. The thickness is X. The wide face YZ is perpendicular to c,, and the faces X Y and X2 are parallel to (IIO),, and (llo),, respectively. The plate edge X and Z are parallel to a,. and co-, respectively, but the edge Y is slightly inclined to b.-. The length of the carbide image in the direction perpendicular to a,. is less than 30 A after tempering for 1 day, about 70 A after 20 days and about 100 A after 100 days. On the other hand, the plate thickness is 30-50 A and almost independent of the period of tempering. After tempering for 100 days, the spacing D is about half of the plate thickness. Since the carbide is thinnest in the a,” direction, long streaks might be expected to run along the [lOO],-* direction in the diffraction patterns. Actually, however, the streaks were never observed in this direction. The disappearance of the streaks can be explained as follows. For simplicity, let us consider the case where the v-carbide plates are arranged in parallel with the a,. axis. A simple calculation gives the diffraction intensity from this scattering system I(H) = ]P(H)]2. L,

. L, - L,

= sin2 ?TN{(X + D)/A)H * A

L A

sin2 r{ (X + D)/A}H * A

sin2 n(X/A)H

VOL.

20,

1972

tml

\I

-A Q.”

FIG. 17. Arrangement of the carbide plates in the matrix. The carbide plates are aligned along a dislocation line.

Fig. 18. It becomes very sharp in comparison with the Laue function of one plate L,. In the above, the carbide plates are assumed to have the same thickness and to be arranged exactly parallel with constant interval. However, it is not necessary for these conditions to be strictly adhered to. The sharpening of the diffraction spots is caused by the remarkable interference among the diffracted waves from the aligned plates. In fact, the images of the aligned carbide plates are slightly different in contrast as

(4) ’A

sin2 ~TH* A

’

(5) L

= sin2 W(Y/B)H B

.B (6)

sin2 7rH B l

L

A = c,,

c

= sin2 r(Z/C)H

l

C

sin2 ~TH* C

B = a, + b,,

C = -a,, + b,

(7) (8)

A, B and C are lattice vectors along the edges X, Y and 2, respectively, H the scattering vector, F(H) the crystal structure factor and N the number of the carbide plates under the same Bragg condition. From (5), we can see that the Laue function of one plate is modified by the periodic array of the plates. For the case of X/A = 10, D/A = 5 and N = 4, the Laue function along the [loo],.* axis, LA, is shown in

FIQ. 18. Lam functions along [lOO]cr-*axis. L

A

+ D)/A}H *A. sin* n(X/A)H me a{(X + D)/A}E .A sin*&-A L = h* +X/A)E - A 1 sin*wH-A ’

= siy* n N{(X

Curves are drawn for the case X/A N = 4.

-A ’

= 10, D/A = 5 arId

HIROTSU

A?LD

NAGAKURA:

CARBIDE

PRECIPITATED

DURING

TEMPERING

653

0.5 w Cal

(b 1

(cl

FIG. 19. (a) Bright field image of a specimen tempered for 1 da (b) corresponding diffraction pattern and (c) dark field image using q-P01 spot. seen in Fig, 15(c), indicating a slight mutual inclination between the plates. The carbide plates are small in directions perpendicular to the aa” direction. This may cause streaks in the diffraction patterns. Detailed examination of various electron diffraction patterns revealed that all the reciprocal lattice points have spikes along the [llO],* and the [TlO],* directions, in other words, spikes along the [OlO],.* and a direction slightly inclined to the [OOl],.* direction, in specimens tempered less than 20 days. In the case of tempering for 100 days, the spikes exist only in the [ilO],* direction. Figure 19(a) is a bright field image of a specimen tempered for 1 day, (b) the corresponding diffraction pattern and (c) the dark field image taken with q-i01 spot. In the images, we can observe dotted lines due to aligned carbide particles. The image width is less than 30 A in the c,. direction projected onto the specimen surface. Corresponding to this, long streaks are observed in (b) to run along the [210],* direction, which coincides with the [ilO],* direction projected along the incident beam direction. In Fig. 15(b), we can also see short [IlO],* streaks running along the [OOl],** direction. Figures 11(a)-(c) show the change of streak length along the [IlO],* direction with increasing period of tempering. In (a) for 1 day tempering, the streaks are so long as to bridge the adjacent diffraction spots and in (b) for 20 days tempering they are shortened, but in (c) for 100 days tempering they become too short to be observed. In Fig. 6(b), the carbide spots also have streaks along the [llO],* or [OlO],.* direction. These facts show the growth of the carbide particles in the direction perpendicular to the (IlO), and (llO), planes with

the period of tempering. The streaks cannot be ascribed to stacking disorder of the (TlO), and (llO), planes, because the q-carbide structure does not allow stacking fault formation on these planes. In the matrix lattice, 8,” and b,. are crystallographitally equivalent. Accordingly, the carbide particle can grow with either the orientation relationship (3) or (llO), )I (loo),. and [OOl], ]I [OlO],.. If individual carbide particles take the different orientation relationships, the diffraction patterns may contain two sets of the carbide net patterns. However, not many cases gave such patterns, indicating that the carbide particles prefer to take one of the two possible orientation relationships in a limited area. This is somewhat different from the observation by Wells,“i) who reported that the carbide frequently precipitated on more than one plane of the martensite lattice. 4. DISCUSSION

We have shown that the carbide precipitated during the first stage of tempering of martensitic high carbon steel is v-carbide or g-Fe&. This is orthorhombic and isomorphous with Co,N and Co&. This result differs from that obtained by Jack@) from X-ray diffraction. He found s-carbide under almost the same specimen conditions. X-ray powder diffraction lines are very weak and inappropriate for detailed structure analysis. In fact, our X-ray powder patterns could be interpreted as either q or e-carbide. The difficulty in the X-ray investigation cannot be avoided by the use of single crystal specimens, since martensitic steels of high carbon content are twinned and the precipitated carbide is too fine. In this study, it was shown that precise electron micrographs and diffraction patterns

654

ACTA

METALLURGICA,

VOL.

20,

1972

are effective for the determination of the crystal structure of fine precipitates. Accordingly, it may be necessary to r-e-examine the conclusion that the precipitated carbide in Fe-K-C alloys was s-carbide. Our conclusion does not deny the presence of Ecarbide. s-carbide and q-carbide are different carbides and the former has been synthesized by several authors.‘7,s) There is a structural similarity between the two. s-carbide has a hexagonal unit cell with dimensions a, = 2.752 and c, = 4.353 A,(l’) which can be transformed to an orthorhombic cell with dimensions a,’ = &a, = 4.767, b,’ = c, = 4.353 and ce’ = a, = 2.752 A. Comparing these values with the lattice parameters of q-carbide, we obtain a, = (1 - O.O13)a,‘, b, = (1 - O.OOS)b,’ and c,, = (1 + O.O28)c,‘. Thus, the distance between the closed packed layers of iron atoms does not change appreciably, but the iron atoms in the q-carbide aredisplaced especially severely in c,’ or a, direction. The moat remarkable difference lies in the carbon atom arrangement. In the s-carbide, the carbon atom arrangement has hexagonal symmetry, but in the q-carbide, it ha8 orthorhombic symmetry. According to the orientation relationship described in Section 3.3, the arrangements of iron atoms on (OOl), and (lOO),. planes are superposed in Fig. 20. Here, the plain and solid circles represent iron atoms in the q-carbide and in the matrix a”, respectively, and the corresponding atoms are connected by arrows. The matching of atoms is fairly good. In Section 3.4, we showed that the carbide particles may be surrounded by (llO), and (lIO), planes as well as the habit plane (OOl),. The external shape projected on t’he (OOl), plane is the same as the cell shape shown by the broken lines in the figure. This cell has the basic vectors A, B and C defined by (8). The side face (llO), is parallel to (OlO),. and the direction [ilO], is parallel to [OOl],.. Therefore, the misfit parameters between the carbide and the matrix in a,., b,. and c,. directions are (c,, - am-)/a,. = -0.007, (do,,,, - b,.)/ b,. = 0.116 and (C/2 - c,-)/c,- = 0.104, respectively. The misfit parameter in the a,. direction is so small that r-carbide can grow initially in this direction. On the other hand, the large misfit parameters in the b, and c,- directions may be responsible for the slow rate of the carbide growth in these directions. This explains the carbide growth with increased period of tempering. In the present study, it was found that the carbide precipitates periodically along dislocations and does not grow beyond a thickness of about 50 A in the a,.‘direction even after 100 days tempering. This phenomenon is very interesting, but its meohanism is not clear at present.

FIG. 20. Arrangement of iron atoms on (001)~ and ( 100)tl* planes: open circles, iron atoms for q-carbide; filled circles, iron atoms for the matrix of tempered martensite. Heavy lines represent the orthorhombic unit cell of v-carbide with lattice vector •~, bq and cq; broken lines indicete the monoclinic cell with lattice vector A, B rmd C.

Our study involved hyper-eutectic carbon steel. In hypo-eutectic steels, especially when the carbon content is less than 0.4 wt. ‘4, the twinned structure is not formed by quenching.(“) It is important to study the structure and morphology of precipitates in this case. Furthermore, it is desirable to make a detailed study of the carbide precipitated from quench-aged a-iron containing carbon less than 0.03 wt. %, although the precipitate has been believed to be e-carbide. These studies are now in progress. ACKNOWLEDGMENTS

The present authors wish to express their thanks to Professor S. Oketani for his guidance throughout this work. Thanks are also due to Professor M. Kikuchi for helpful discussions. REFERENCES 1. G. KURDJUMOVend L. LYSSAK, J. Iron Steel Inat. 166, 29 (1947). 2. C. 8. ROBERTS,B. L. AVEBBACE and M. COHEN, Trans. Am. sot. Mctola 4s. 576 (1953). 3. B. S. LEMENT,B. L. AVERISACR and M. COHEN,Tranu. Am. sot. Metda 47, 291 (1955).

HIROTSU

AND NAGAKURA:

CARBIDE

1. B, S. LEMEIT and itI. COHEX, Acta Met.,& 469 (1956). 5. R. D. HEIDENBEICH, L. S~URKEY &nd H. L. WOODS,J.

appt. Phys. 17, 127 (1946). 6. K. H. JACK, J. Iron &eel Inst. 169.26 (1951). 7. L.J.E. HOFER,E.M.COHN sndW.C. PEEBLES,J.Am. Chem. Sot. 71, 189 (1949). 8. S. OKETANIend S. NAQAKURA, J. Japan Inst. Metals 20, 460 (19561. 9. M. $. A&JZOV and B. V. KHAYENKO, Phy8b &fetal8 ~~~~~~. 15, 128 (1962). 10. E. TEK~Nand P. M. KELLY, P~~p~~ From Iron Base AZby8. Gordon &nd Breach { 1965). 11. M. G. H. WELLS, Ado Met. Ifl, 389 (1964). 12. G. J. BARTON,Acta Met. 17, 1085 (1969).

PRECIPITATED

DURIXG

TEMPERIXG

655

and S. OXETASI, Int. Co@. Sci. 13. Y. Hrxupsr;, S. ?~‘A~AKVRA Tech. fron and SteeE, ,Tokyo (1970); Trans. Iron SfeeE Inet. Japan 11,Suppt. 1140 (1971). 45 14. P. M. KELLY and J. NUTTISG, Proc. R. Sot. MB, (1900). IL Sarnmv. J. cryatallogr. Sot. Japan. 11,SO (1969). ::: R. L. PATTERSOXand C. M. WAYMAS, Acta Met. 14, 347 I19661. 17. 6. NA&AX~XA, J. Pkys. Sac. Jupm 14,186 (1959). and S. ORETAXI, Trans. iron SteeEInst. 18. S. NAOAKZ~I~A Japan 8,265 (1968). 19. W. PITSCRand A. SCRRADER.drck. E~e~kjit~~e~. 29,715

(1968).