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Carbides in low-temperature-carburized stainless steels
Acta Materialia 52 (2004) 1469–1477 www.actamat-journals.com
Carbides in low-temperature-carburized stainless steels F. Ernst *, Y. Cao, G.M. Michal Department of Materials Science and Engineering, Case Western Reserve University, 10900 Euclid Avenue, Cleveland, OH 44106-7204, USA Received 12 July 2003; received in revised form 24 November 2003; accepted 26 November 2003
Abstract A novel, low-temperature (470 °C) gas-phase carburization treatment, developed by the Swagelok Company, increases the surface hardness of 316-type austenitic stainless steels from 200 to 1000HV25 via a colossal supersaturation of up to 12 at.% carbon in solid solution. Upon extended treatment, carbide precipitation does eventually occur. Transmission electron microscopy of carburized bulk, foil, and powder samples has revealed two different carbides: M5 C2 (‘‘H€ agg’’ or ‘‘v’’ carbide) and M7 C3 . The particles of both carbides have the shape of long needles or laths, contain a high density of planar defects normal to the particle axis, and adopt special topotaxial orientation relationships with the austenite matrix. Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Austenitic steels; Surface alloying; Colossal supersaturation; Carbides; Electron diffraction
1. Introduction In a previous publication [1], we have reported that a novel, low-temperature (470 °C) gas-phase carburization treatment, developed by the Swagelok Company, increases the surface hardness of 316-type austenitic stainless steels from 200 to 1000HV25. In addition, the treatment improves the corrosion resistance. While normally the precipitation of carbides restricts the carbon concentration in the austenite of 316-type steels to <0.015 at.% at 470 °C, the Swagelok treatment generates a colossal supersaturation of up to 12 at.% carbon in solid solution – 800 times the equilibrium solubility at the carburization temperature (thus an even larger supersaturation at room temperature). Upon extended treatment, however, carbide precipitation does eventually occur. By comparing experimental X-ray diffractograms of low-temperature carburized powders of 316-type austenitic stainless steels with simulated diffractograms of various carbides, we have concluded that the carbide precipitates have the structure of Fe5 C2 , known as ‘‘H€ agg’’ or ‘‘v’’ carbide (space group: C2/c, Table 3) [1]. *
Corresponding author. Tel.: +1-216-368-0611; fax: +1-216-3683209. E-mail address: [email protected] (F. Ernst).
Here, we report transmission electron microscopy (TEM) studies of the carbides forming in excessively low-temperature-carburized bulk, foil, and powder specimens of several different 316-type stainless steels.
2. Experimental methods Table 1 lists the composition of the materials we have studied by TEM (in at.%, prior to carburization, major elements only). Table 2 indicates details of the carburization treatment for each type of specimen: the annealing treatment prior to carburization, the carburization temperature TC , and the carburization time tC . All specimens were provided in the carburized state by the Swagelok company using repeated process runs, adding up to the carburization times listed in Table 2. For preparing plan-view and cross-sectional TEM specimens from the bulk samples KMD-A and KJL-A, we employed standard techniques, including backthinning and thinning of sandwich slices, dimple grinding, and Arþ ion-beam milling with a Gatan model 691 precision ion polishing system (PIPSTM ). The primary energy of the Arþ ions was 4 keV and the beam current was kept below 25 lA. Plan-view specimens of the carburized foils ADT15 and ADT40 were also prepared by
1359-6454/$30.00 Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2003.11.027
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Table 1 Major elements of the studied materials (at.%)
KMD-A KJL-A ADT40 ADT15 P42860
C
Mn
Si
P
S
Cr
Ni
Mo
Fe
0.23 0.21 0.16 0.16 0.11
1.79 1.63 1.85 1.61 1.35
1.18 0.34 1.18 0.98 0.79
0.04 0.03 0.05 0.05 0.04
0.04 0.05 0.01 0.01 0.01
18.45 18.50 17.30 18.76 17.97
11.62 12.55 9.53 10.65 9.85
1.19 1.51 1.23 1.10 1.22
64.94 64.68 67.69 65.29 68.65
Table 2 Carburization treatments of the studied materials Sample
Morphology
Annealed
TC (°C)
tC (h)
KMD-A KJL-A ADT40 ADT15 P42860
Bar stock Bar stock 100-lm thick foil 38-lm thick foil Powder, £ 16 lm
1.0 h at 1065 °C 1.0 h at 1065 °C 1.5 h at 1065 °C 1.5 h at 1065 °C –
470 470 465 465 465
246 246 44 44 149
the standard back-thinning technique. In order to make cross-sectional TEM specimens from the foils ADT15 and ADT40, we applied the method developed by Strecker et al. [2]: a stack of foil sections was mounted in the slit of a cylindrical alumina rod, which was then inserted in an alumina cylinder with an outer diameter of 3 mm. After fixing the assembly with M-bond 600e adhesive, 1 we cut slices from the cylinder and prepared these for TEM by the standard techniques described before. The powder sample P42860, finally, was prepared for TEM by dispersing it in methanol and depositing a drop of the dispersion onto a holey carbon film, supported by a copper grid. Conventional bright-field images and electron diffraction patterns of all specimens were obtained with a Philips CM20 microscope, operated at an accelerating voltage of 200 kV, and a JEM4000EX (JEOL), operated at an accelerating voltage of 380 kV. Both instruments were equipped with a LaB6 thermal emitter. For simulating electron diffraction patterns, and for drawing crystallographic ball models of the carbide structures, we employed the software package JEMS [3], the JAVA version of EMS [4] by Stadelmann, as well as the software package EMS Online [5].
3. Results Fig. 1 presents a conventional bright-field TEM image of a cross-sectional specimen of the foil ADT15. The micrograph reveals heavily dislocated austenite (dark regions), labeled c, intermixed with needles (or laths) of a second phase, labeled v, about 50 nm in diameter and up to several micrometers long. As shown in the fol-
1
Vishay micro-measurements.
Fig. 1. Cross-sectional TEM bright-field image of carbide particles in the carburized foil ADT15.
lowing, these particles are v-carbide. Often, the needles lie perpendicular to the surface. As seen in Fig. 1, they can be present with a higher volume fraction than the austenite. The straight vertical features normal to the needle axes in Fig. 1 indicate the presence of a high concentration of planar faults in the needles. TEM images of KMD-A, KJL-A, and ADT40 reveal similar scenarios: heavily dislocated austenite, containing elongated particles of a second phase, invariably containing a high density of stacking faults or microtwins normal to the long axis. According to the results we obtained by XRD (X-ray diffraction) [1], the carbide particles in Fig. 1 should have the structure of Fe5 C2 (H€agg carbide or v-phase, Table 3), shown in Fig. 2. Attempts to study the crystal structure of the carbide particles by convergent-beam electron diffraction (CBED) remained unsuccessful because the patterns did not reveal any HOLZ lines, or
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Table 3 Crystal structures Carbide
Lattice parameters
Space group
Atom positions
v-Fe5 C2 (H€ agg carbide)
a ¼ 1:1552 nm, b ¼ 0:4564 nm, c ¼ 0:5043 nm, a ¼ c ¼ 90° b ¼ 97:68° obtained by fitting experimental data with Celref software [24]
C2/c (15) monoclinic
Fe in 0.097, 0.078, 0.423; Fe in 0.215, 0.581, 0.306; Fe in 0, 0.561, 0.25; C in 0.107, 0.285, 0.149 obtained from [25]
Fe7 C3 a
a ¼ 0:4537 nm, b ¼ 0:6892 nm, c ¼ 1:1913 nm a ¼ b ¼ c ¼ 90° (JCPDS file 75-1499)
Pnma (62) orthorhombic
Fe in 0.25, 0.07, 0.02; Fe in 0, 0.07, 0.81; Fe in 0.25, 0.25, 0.20; Fe in 0.25, 0.25, 0.42; Fe in 0, 0.25, 0.63; C in 0.11, 0.03, 0.35; C in 0.38, 0.25, 0.57 obtained from [15]
a
Fe7 C3 is isostructural to Cr7 C3 with a ¼ 0:4532 nm, b ¼ 0:7015 nm, c ¼ 1:2153 nm (JCPDS file 36-1482).
Fig. 2. Crystallographic model of Fe5 C2 , the ‘‘H€agg’’ carbide (four unit cells in ½0 1 0 projection).
other three-dimensional information, and the pattern in the discs was too irregular for symmetry considerations – apparently as a result of the high defect density and internal stresses. However, Fig. 3(a), the selected-area electron diffraction pattern of the region shown in Fig. 1, confirms the hypothesis that the carbide particles in Fig. 1 have the structure of Fe5 C2 : According to the simulated diffraction pattern in Fig. 3(b), Fig. 3(a) corresponds to the pattern expected for H€ agg carbide in ½0 1 0 projection. The fine streaks at each spot in Fig. 3(a) apparently originate from the planar faults (stacking faults or twins) normal to the long axes of the carbide needles; our rotation calibration for the CM20 indicates that they are oriented orthogonal to the fault plane. The orientation of the streaks in Fig. 3(a) indicates that the planar faults are parallel to the plane spanned by the unit cell vectors b and c of the monoclinic crystal structure. Accordingly, the long axis of the needles corresponds to the a axis of the (monoclinic) unit cell. The diffraction pattern of Fig. 3(a) also reveals faint reflections from the austenite matrix. According to these reflections, the austenite had a h2 1 1i axis parallel to the incident electron beam, and the following orientation relationship exists between the carbide (index v) and the austenite (index c): h0 1 0iv k h2 1 1ic ; f0 0 1gv k f1 1 1gc :
ð1Þ
The selected-area diffraction pattern of Fig. 3(c), recorded in a different TEM specimen (a cross-sectional specimen of ADT15), reveals the same common diffraction pattern of carbide and austenite as Fig. 3(a). This indicates that the orientation relationship (1) is significant. In fact, this orientation relationship seems to be unique; evaluating the diffraction patterns of about 20 further H€agg carbide particles, we have not observed any other orientation relationship. The particles of Fig. 1 are all within the same austenite grain, and they are all aligned with respect to each other. This means that not only their lattice has a preferred orientation relationship with the austenite lattice, but also that their long axis corresponds to a particular crystallographic direction in the carbide (b c) and in the austenite. Elongated carbide particles, similar to those in Fig. 1, were also observed in plan-view and cross-sectional specimens of carburized KMD-A, KJL-A, and ADT40, the thicker one of the two foil samples. As an example, Fig. 4(a) shows a plan-view dark-field TEM image of carbide particles in a cross-sectional specimen of KMDA. The corresponding selected-area electron diffraction pattern in Fig. 4(b) is consistent with the hypothesis that these particles have the same crystal structure as those in Fig. 1, thus the structure of Fe5 C2 . The pattern corresponds to that of Fe5 C2 in ½0 0 1 projection, and the reflection that served for recording the image Fig. 4(a) is encircled and identified as a ð5 3 0Þ reflection. Not only the structure, but also the morphology of the carbide particles, which show up bright in the dark-field image of Fig. 4(a), closely corresponds to those in ADT15 (Fig. 1). Evaluation of about 50 electron diffraction patterns of needle-shaped carbide particles in all bulk and foil samples agree with the hypothesis that most of these particles have the crystal structure of H€agg carbide. Diffraction patterns consistent with the structure of Fe5 C2 are also observed in TEM diffraction patterns from the powder specimen P42860. According to the results of our XRD studies, reported in [1], the powder is fully converted to carbide. Our TEM studies have confirmed this finding. Fig. 5 shows two examples of small (electron-transparent) powder particles and the corresponding selected-area diffraction patterns. These
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Fig. 4. Carbide particles in KMD-A. (a) Plan-view dark-field TEM image. (b) Corresponding selected-area electron diffraction pattern. The ð5 3 0Þ reflection that served for recording the image (a) is encircled.
Fig. 3. Selected-area diffraction patterns. (a) Pattern recorded in the region shown in Fig. 1, rotated relative to Fig. 1 according to our calibration of image rotation. (b) Simulated diffraction pattern of H€ agg carbide in ½0 1 0 projection. (c) Pattern recorded in a different specimen, indicating the same orientation relationship between the H€ agg carbide and the austenite (c) as in (a).
patterns, recorded in ½1 0 0v and ½1 1 0v direction, exclusively exhibit carbide reflections. The powder particles possess pronounced facets, corresponding to lowindexed crystallographic planes, and the entire morphology – particularly the special angles the facets make with each other – reflects the low symmetry of the monoclinic crystal structure.
In summary, our observations lead to the conclusion that carbide particles with the crystal structure of Fe5 C2 (H€agg carbide) have formed in all five carburized specimens. Since 316-type steels also contain major amounts of Cr and Ni, it seems likely that these elements are also present in the carbide phase. Indeed, high-resolution analytical TEM has confirmed this hypothesis [6]. Based on the results from XRD, the observation of H€agg carbide by TEM was expected. However, we also observed particles of another carbide, which had not been detected by XRD before. The electron diffraction patterns of this carbide are consistent with the crystal structure of Fe7 C3 (Table 3, Fig. 6). Fig. 7 presents a selected-area electron diffraction (SAD) pattern of such a carbide particle in a plan-view, back-thinned specimen of the foil sample ADT40. This pattern exhibits sixfold symmetry, indicating the existence of either a threefold or a sixfold symmetry axis in the crystal structure. Since other diffraction patterns we have obtained from these carbide particles, including the pattern of Fig. 9 discussed below, are incompatible with cubic symmetry, Fig. 7 suggests that this carbide has either hexagonal or trigonal/rhombohedral symmetry. In both cases, the
F. Ernst et al. / Acta Materialia 52 (2004) 1469–1477
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Fig. 7. Selected-area diffraction pattern of a carbide particle in the foil sample ADT40. The inset at the lower left corner shows an enlargement of the area in the small dashed square above. The circled reflections originate from a neighboring austenite grain.
Fig. 5. TEM bright-field images of two powder particles from the carburized powder sample P42860 and corresponding selected-area diffraction patterns. The labels in the diffraction patterns refer to the encircled reflections.
Fig. 6. Crystallographic model of Fe7 C3 (four unit cells in [1 0 0] projection).
structure can be described by a hexagonal Bravais lattice with lattice parameters aC and cC and the viewing direction of Fig. 7 corresponds to the ½0 0 0 1 or ½0 0 0 1 direction of the hexagonal crystal lattice. Fig. 8 presents an example for the particle morphology of second carbide, observed in a cross-sectional TEM specimen of KJL-A. The appearance of these particles closely resembles that of the M5 C2 carbide particles of Fig. 1: They are elongated, needle-shaped, with a diameter of a few tens of a nanometer, and
Fig. 8. Cross-sectional TEM bright-field image of carbide particles the carburized foil ADT15. These particles have a different crystal structure than those of Fig. 1.
contain a high concentration of planar faults (particularly obvious in the arrowed particle, accidentally oriented with the fault planes precisely parallel to the viewing direction). Correlating TEM images and diffraction information (using the image rotation calibration of our instrument) revealed that the ½0 0 0 1 direction of the hexagonal axis coincides with the long axis of the carbide grains of Fig. 8. The six encircled, weak diffraction spots in Fig. 7 are f2 2 0g reflections of adjacent austenite, which was also included in the SAD aperture. Accordingly, the main symmetry axis of the carbide lies parallel to a h1 1 1i
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direction of the austenite. The h1 1 0i directions of the austenite, however, do not coincide with prominent, low-indexed directions in the carbide. As can be seen in Fig. 7 – most clearly in the enlarged inset at the bottom left – each f2 2 0g austenite reflection nearly coincides with a f5 4 1 0g reflection of the carbide (a reflection of the type 5aC þ bC in Fig. 7). It appears that the 3% misfit between the corresponding lattice planes in the carbide and in austenite prevents the carbide from aligning its lattice with the more densely packed f2 2 0g planes of austenite. Consequently, the translations aC and bC in the carbide pattern of Fig. 7 approximately correspond to h5 4 1i directions of the austenite, not h1 1 0i. By evaluating the carbide f1 1 0 0g spacing, corresponding to aC in Fig. 7, relative to the f2 2 0g spacing of the austenite, and by measuring the expanded lattice parameter at the surface of ADT40 by XRD, which yielded ac ¼ 0:3661 nm, we obtained kaC k
1
¼ ð0:616 0:005Þ nm:
ð2Þ
For the lattice parameter a of the carbide phase parallel to the basal-plane, this implies 2 1 aC pffiffiffi kaC k ¼ ð0:711 0:006Þ nm: 3
ð3Þ
(The error limits given here and in the following exclusively reflect the error in determining the positions of the diffraction spot intensity maxima.) Fig. 9 shows another diffraction pattern obtained from a carbide particle of the second kind. This pattern shows pronounced streaks, reflecting the high density of planar faults observed in Fig. 8, and exhibits mm symmetry. Like Fig. 7, the pattern includes diffraction spots from the carbide phase as well as from the neighboring austenite. Since the pattern includes f1 1 1g as well as
Fig. 9. Selected-area electron diffraction pattern of an M7 C3 carbide particle in h2 1 1 0i projection. The bold reflections originate from a neighboring austenite grain in h3 2 1i projection.
f3 3 1g reflections of austenite, the viewing direction corresponds to h3 2 1i in the austenite. The lattice planes producing the fundamental diffraction vector aC lie parallel to f5 4 1g planes in austenite. Measuring the length of aC in Fig. 9 relative to the distances of the austenite spots from 000 and calibrating the austenite spacings with the lattice parameter ac ¼ 0:3661 nm we obtained by XRD for the expanded austenite at the surface of ADT40, we obtained kaC k
1
¼ ð0:613 0:002Þ nm:
ð4Þ
This result deviates by less than 0.5% from the result (2), and therefore identifies the reflecting planes as f1 1 0 0g planes of the hexagonal lattice. The lattice planes producing the fundamental diffraction vector cC lie parallel to f1 1 1g planes in the austenite. According to the information obtained from Fig. 7 and in agreement with the 90° angle they make with the f1 1 0 0g planes, the planes causing the reflection cC constitute the basal planes of the carbide. Hence, the direction in which the pattern of Fig. 9 was obtained corresponds to a ½2 1 1 0 direction of the hexagonal lattice. For the plane spacing that corresponds to cC we obtained kcC k
1
¼ ð0:455 0:002Þ nm:
ð5Þ
At this point, we can justify our assumption that the structure of the second kind of carbide is either hexagonal or rhombohedral, but not cubic. Consider the ratio of the plane spacings (4) and (5) q :¼
kcC k1 kaC k
1
¼ ð0:739 0:004Þ
ð6Þ
If the crystal structure was cubic, Fig. 7 would be the diffraction pattern for a h1 1 1i direction of the primary electron beam, and the fundamental reflections (corresponding to aC ) would be of the type fh h 0g. The pattern of Fig. 9 would correspond to a primary beam direction orthogonal to one of these reflections, for example orthogonal to ½h h 0. Then, the reflections corresponding to cC must have the form ðk k lÞ. Usually, h ¼ 1 for SC (single cubic) or BCC (body-centered cubic) structures and h ¼ 2 for a FCC (face-centered cubic) structures. However, there are no sets of planes in SC/BCC and FCC crystals that would yield a q fulfilling the condition (6). The closest match is obtained for ðk k lÞ ¼ f0 0 2g in SC/BCC and ðk k lÞ ¼ f0 0 4g in FCC, yielding q ¼ 0:707 in each case, but the discrepancy with our result (6) clearly exceeds the error limits. It may be argued, of course, that the 1 1 0 reflections (in SC/BCC) or 2 2 0 reflections (in FCC) could be kinematically forbidden. Under this assumption, it is indeed possible to find plane spacings that match with our result (6). For example, if the fundamental reflections in Fig. 7 were not f2 2 0g but f4 4 0g reflections of a hypothetical FCC structure, the reflections corresponding to cC could be of
F. Ernst et al. / Acta Materialia 52 (2004) 1469–1477
the type f5 5 3g, yielding q ¼ 0:736. However, this would imply not only a primary beam direction with very high indices, h3 3 1 0i in this case, but also a very large lattice parameter: 3.47 nm according to (4) and (5). Considering that Cr23 C6 , for example, has a lattice parameter of 1.07 nm and 116 atoms in the unit cell, this would mean that there are about 4000 atoms in the unit cell. While not impossible, this is at least highly unlikely. Assuming further extinctions in the h1 1 1i zone lead to even more unrealistic results. Therefore, we have discarded the possibility that the observed carbide of the second type has a cubic structure. The diffraction pattern of Fig. 10, obtained from another cross-sectional TEM specimen of ADT40, confirms the lattice parameters and the special orientation relationship deduced from the patterns in Figs. 7 and 9. Summarizing, the crystallites of the second carbide phase also adopt a special orientation relationship with the residual austenite h0 0 0 1iC k h1 1 1ic ; h1 1 0 0i k h5 4 1i ; C
ð7Þ
c
which implies h2 1 1 0iC k h3 2 1ic :
ð8Þ
While the close-packed planes of both phases are parallel, the close-packed directions are not, and there are no common mirror planes of the two adjacent crystal structures parallel to the h0 0 0 1iC k h1 1 1ic direction. According to the diffraction pattern in Fig. 7, the orientation relationship (7) optimizes the alignment of f5 4 1 0gC planes with f2 2 0gc planes, which have nearly the same spacing.
Fig. 10. Selected-area electron diffraction pattern of an M7 C3 carbide particle in h2 1 1 0i projection, similar to Fig. 9, but recorded in a different specimen. Like in Fig. 9, the bold reflections originate from a neighboring austenite grain in h3 2 1i projection, confirming the significance of the corresponding orientation relationship.
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Among the relevant iron carbides, the lattice parameters (2) and (5) only match well (within the limits of error) with those of Fe7 C3 as given by Herbstein and Snyman [7–9]. While the XRD data of those authors suggested a hexagonal structure with a ¼ 0:6882 nm;
c ¼ 0:454 nm;
ð9Þ
later work of Fruchart and Rouaoult [7] revealed that this carbide actually has orthorhombic crystal symmetry (Fig. 6) with the lattice parameters a ¼ 0:454 nm;
b ¼ 0:6879 nm;
c ¼ 1:1942 nm;
ð10Þ pffiffiffi implying c=b 3. This structure was confirmed by TEM studies of Morniroli et al. [10–13]. These authors found that M7 C3 carbides, including Fe7 C3 , Cr7 C3 , and Mn7 C3 , tend to contain a high density of planar defects, mainly twin and antiphase boundaries, between the three different orientation variants of the orthorhombic unit cell. It is because of these defects that the material appears to have hexagonal crystal symmetry under the typical experimental conditions of electron diffraction in a transmission electron microscope. A high density of planar defects seems to be an intrinsic feature of Fe7 C3 and was in fact already reported by Dyson and Andrews in 1969 [14]. With reference to the true, orthorhombic crystal lattice, the viewing direction of Fig. 7 corresponds to ½1 0 0 and the viewing direction of Figs. 9 and 10 to ½0 1 0. While the lattice parameters of Fe7 C3 , Cr7 C3 , and Mn7 C3 [15] are too similar to be safely distinguished under our experimental conditions. Fig. 8 demonstrates that locally the volume fraction of the M7 C3 carbide can be as high as 75%. It seems unlikely, therefore, that this carbide is devoid of iron. On the other hand, analytical TEM has revealed that the carbide particles are not pure Fe7 C3 , either. As will be detailed elsewhere [6], X-ray
Fig. 11. Selected-area electron diffraction pattern of carbide particles in KJL-A.
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energy-dispersive spectroscopy indicates that the major cation species is iron, but the carbides also include significant concentrations of the substitutional alloying elements (Cr, Ni). Another example for the general formation of carbide particles consistent with the Fe7 C3 structure is shown in Fig. 11, which presents a selected-area diffraction pattern of carbide particles in KJL-A. This diffraction pattern corresponds to the structure of Fe7 C3 viewed in ½2 1 1 direction.
4. Discussion To our knowledge, carbide formation in low-temperature-carburized austenitic steels has not been observed before, and the formation of H€ agg carbide (M5 C2 ) in austenitic stainless steel is very unusual. A substantial database exists, on the other hand, about the formation of M5 C2 and M7 C3 carbide under other experimental conditions. In ferritic steels, Dirand and Alqier [16] have observed Fe5 C2 as part of a transformation sequence that leads from martensite to a phase equilibrium of ferrite (a-Fe) and cementite (h-Fe3 C) [16]. In this sequence, v-Fe5 C2 forms from g-Fe2 C and subsequently decomposes into h-Fe3 C. The transformation v ! h occurs progressively through variations in the periodicity of twinning, while the slight change in stoichiometry is apparently accommodated by incorporating planar defects. K€ oninger et al. [17] and Hammerl et al. [18] have observed exactly the opposite sequence of carbides on tempering carbon-rich Fe–C solid solutions prepared by ion implantation. The reason for the observed transformation sequences and the topotaxial orientation relationships (‘‘lattice-invariant deformation’’) between subsequent carbides is seen in the structural similarity between the carbide phases g, v, and h: the carbon atoms reside in triangular prisms of iron atoms, and the individual carbide structures just slightly differ by the patterns in which these prisms are arranged. Given the structural similarity between g; v and h carbide, we speculate that annealing of the low-temperature-carburized material will readily lead to the formation of h (M3 C). On the other hand, the observations reported in the literature raise the question whether the M5 C2 carbide we observe perhaps forms via a precursor of g carbide (M2 C). From our experimental data, however, there is no evidence for the formation of g carbide in low-temperature carburized austenitic stainless steels. M7 C3 is unique in that it has the highest carbon content among the carbides considered here. Also this carbide is built up from structural units of carbon atoms near the center of trigonal prisms of Fe atoms [19]. At present, it is not clear whether what controls the ratio of
M7 C3 to M5 C2 in our material. It seems possible that M5 C2 actually transforms to M7 C3 , not M3 C, under the far-from-equilibrium conditions of low-temperature carburization. In this case, the volume fraction of M7 C3 , which is so much smaller than that of M5 C2 , should increase with increasing time. Tempering experiments are being carried out in our group to further investigate this point. The observation that planar faults in the M5 C2 carbide occur on the planes spanned by the unit cell vectors b and c agrees well with the building principle of the structure discussed by Jack and Wild [19], although we believe that the defect structure reported in the present publication has not been observed before. The formation of these faults may be related to transformation stresses or to positioning errors on the growth face. In M7 C3 carbides, however, a high density of planar defects has also been observed in single crystals – absent transformation-related stresses [10,12,20]. For the carbide particles of both phases, M5 C2 and M7 C3 , preferred orientation relationships exist with the austenite, and the long axis of both types of carbide particles is aligned with a particular crystallographic direction. Crystallographic orientation relationships have often been interpreted in terms of minimization of lattice mismatch between the adjacent phases. Whether this interpretation is correct for the extreme non-equilibrium conditions of carbide formation in austenite with a colossal supersaturation of carbon will be subject of future studies, for example by analyzing the observed orientation relationships with the model of Pirouz et al. [21–23]. Another reason for the establishment of a unique orientation relationship could be the atomistic mechanism of carbide nucleation, as discussed for Fe7 C3 by Dyson and Andrews [14], for example a diffusion-less nucleation mechanism involving shear. Concerning the properties of low-temperature-carburized austenitic stainless steel, the formation of carbides is generally undesirable. While carbides are known to increase the flow stress of austenitic stainless steels, our previous results [1] demonstrate that precipitation hardening contributes only marginally to the surface hardness of low-temperature-carburized 316-type austenitic stainless steels. Moreover, since the carbide particles have different mechanical and electrochemical properties than the austenite matrix, their presence may compromise the fatigue and corrosion behavior. The micromechanism of carbide formation is a subject of ongoing research in our group.
5. Conclusions Our TEM studies austenitic stainless steels carburized at low temperature have provided important insight in the microstructural changes that accompany the carbu-
F. Ernst et al. / Acta Materialia 52 (2004) 1469–1477
rization process after prolonged treatment. In addition to the colossal supersaturation of the austenite with carbon and the formation of an M5 C2 carbide, which we have already observed by XRD [1], TEM has revealed a small fraction of a second carbide phase, M7 C3 , and a high dislocation density in the austenite matrix. We conclude that a quantitative understanding of the surface properties resulting from low-temperature carburization must take into account the entirety of all these features. The formation of carbides under the far-fromequilibrium conditions of colossal supersaturation with carbon constitutes a very unusual scenario, which warrants further analysis.
[4] [5] [6] [7] [8] [9] [10]
Acknowledgements
[18]
We thank Peter Williams, Sunniva Collins, and Steven Marx for numerous fruitful discussions, Arthur Heuer and Fumiyasu Oba for comments on the manuscript, and the Swagelok Company for financial support.
[19] [20]
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Stadelmann PA. Ultramicroscopy 1987;21:131. Available from: http://cimesg1.epfl.ch/CIOL/ems.html. Ernst F, Cao Y, Oba F, Michal G [in preparation]. Fruchart R, Rouaoult A. Ann Chim 1969;4:143. Eckstrom, Adcock. J Am Chem Soc 1950;72:1042. Herbstein, Snyman. Inorg Chem 1964;3:894. Morniroli JP, Bauer-Grosse E, Gantois M. Philos Mag A 1983;48:311. Morniroli JP, Gantois M. J Appl Crystallogr 1983;16:1. Morniroli JP, Khachfi M, Courtois A, Gantois M, Mahy J, Van Dyck D, et al. Philos Mag A 1987;56:93. Bauer-Grosse E, Morniroli JP, Le Caer G, Frantz C. Acta Metall 1981;29:1983. Dyson DJ, Andrews KW. J Iron Steel Inst 1969;207:208. Bouchaud JP. Ann Chim 1967;2:353. Dirand M, Afqir L. Acta Metall 1983;31:1089. K€ oniger A, Hammerl C, Zeitler M, Rauschenbach B. Phys Rev B 1997;55:8143. Hammerl C, K€ oniger A, Rauschenbach B. J Mater Res 1998;13:2614. Jack KH, Wild S. Nature 1966;212:248. Morniroli JP, Khachfi M, Gantois M. J Microscopie Spectroscopie Electroniques 1986;11:115. Poulat S, Ernst F. Mater Sci Eng A 2002;323:9. Pirouz P, Ernst F, Ikuhara Y. Diffusion Defect Data B 1998;59– 60:51. Pirouz P, Ikuhara Y, Ernst F. On epitaxy and orientation relationship in physisorption. Interface science and materials interconnection. Tokyo: The Japan Institute of Metals; 1996. p. 107. Available from: http://www.inpg.fr/LMGP/, http://www.ccp14.ac. uk/tutorial/LMGP. Retief JJ. Powder Diffraction 1999;14:130.
Carbides in low-temperature-carburized stainless steels F. Ernst *, Y. Cao, G.M. Michal Department of Materials Science and Engineering, Case Western Reserve University, 10900 Euclid Avenue, Cleveland, OH 44106-7204, USA Received 12 July 2003; received in revised form 24 November 2003; accepted 26 November 2003
Abstract A novel, low-temperature (470 °C) gas-phase carburization treatment, developed by the Swagelok Company, increases the surface hardness of 316-type austenitic stainless steels from 200 to 1000HV25 via a colossal supersaturation of up to 12 at.% carbon in solid solution. Upon extended treatment, carbide precipitation does eventually occur. Transmission electron microscopy of carburized bulk, foil, and powder samples has revealed two different carbides: M5 C2 (‘‘H€ agg’’ or ‘‘v’’ carbide) and M7 C3 . The particles of both carbides have the shape of long needles or laths, contain a high density of planar defects normal to the particle axis, and adopt special topotaxial orientation relationships with the austenite matrix. Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. Keywords: Austenitic steels; Surface alloying; Colossal supersaturation; Carbides; Electron diffraction
1. Introduction In a previous publication [1], we have reported that a novel, low-temperature (470 °C) gas-phase carburization treatment, developed by the Swagelok Company, increases the surface hardness of 316-type austenitic stainless steels from 200 to 1000HV25. In addition, the treatment improves the corrosion resistance. While normally the precipitation of carbides restricts the carbon concentration in the austenite of 316-type steels to <0.015 at.% at 470 °C, the Swagelok treatment generates a colossal supersaturation of up to 12 at.% carbon in solid solution – 800 times the equilibrium solubility at the carburization temperature (thus an even larger supersaturation at room temperature). Upon extended treatment, however, carbide precipitation does eventually occur. By comparing experimental X-ray diffractograms of low-temperature carburized powders of 316-type austenitic stainless steels with simulated diffractograms of various carbides, we have concluded that the carbide precipitates have the structure of Fe5 C2 , known as ‘‘H€ agg’’ or ‘‘v’’ carbide (space group: C2/c, Table 3) [1]. *
Corresponding author. Tel.: +1-216-368-0611; fax: +1-216-3683209. E-mail address: [email protected] (F. Ernst).
Here, we report transmission electron microscopy (TEM) studies of the carbides forming in excessively low-temperature-carburized bulk, foil, and powder specimens of several different 316-type stainless steels.
2. Experimental methods Table 1 lists the composition of the materials we have studied by TEM (in at.%, prior to carburization, major elements only). Table 2 indicates details of the carburization treatment for each type of specimen: the annealing treatment prior to carburization, the carburization temperature TC , and the carburization time tC . All specimens were provided in the carburized state by the Swagelok company using repeated process runs, adding up to the carburization times listed in Table 2. For preparing plan-view and cross-sectional TEM specimens from the bulk samples KMD-A and KJL-A, we employed standard techniques, including backthinning and thinning of sandwich slices, dimple grinding, and Arþ ion-beam milling with a Gatan model 691 precision ion polishing system (PIPSTM ). The primary energy of the Arþ ions was 4 keV and the beam current was kept below 25 lA. Plan-view specimens of the carburized foils ADT15 and ADT40 were also prepared by
1359-6454/$30.00 Ó 2003 Acta Materialia Inc. Published by Elsevier Ltd. All rights reserved. doi:10.1016/j.actamat.2003.11.027
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Table 1 Major elements of the studied materials (at.%)
KMD-A KJL-A ADT40 ADT15 P42860
C
Mn
Si
P
S
Cr
Ni
Mo
Fe
0.23 0.21 0.16 0.16 0.11
1.79 1.63 1.85 1.61 1.35
1.18 0.34 1.18 0.98 0.79
0.04 0.03 0.05 0.05 0.04
0.04 0.05 0.01 0.01 0.01
18.45 18.50 17.30 18.76 17.97
11.62 12.55 9.53 10.65 9.85
1.19 1.51 1.23 1.10 1.22
64.94 64.68 67.69 65.29 68.65
Table 2 Carburization treatments of the studied materials Sample
Morphology
Annealed
TC (°C)
tC (h)
KMD-A KJL-A ADT40 ADT15 P42860
Bar stock Bar stock 100-lm thick foil 38-lm thick foil Powder, £ 16 lm
1.0 h at 1065 °C 1.0 h at 1065 °C 1.5 h at 1065 °C 1.5 h at 1065 °C –
470 470 465 465 465
246 246 44 44 149
the standard back-thinning technique. In order to make cross-sectional TEM specimens from the foils ADT15 and ADT40, we applied the method developed by Strecker et al. [2]: a stack of foil sections was mounted in the slit of a cylindrical alumina rod, which was then inserted in an alumina cylinder with an outer diameter of 3 mm. After fixing the assembly with M-bond 600e adhesive, 1 we cut slices from the cylinder and prepared these for TEM by the standard techniques described before. The powder sample P42860, finally, was prepared for TEM by dispersing it in methanol and depositing a drop of the dispersion onto a holey carbon film, supported by a copper grid. Conventional bright-field images and electron diffraction patterns of all specimens were obtained with a Philips CM20 microscope, operated at an accelerating voltage of 200 kV, and a JEM4000EX (JEOL), operated at an accelerating voltage of 380 kV. Both instruments were equipped with a LaB6 thermal emitter. For simulating electron diffraction patterns, and for drawing crystallographic ball models of the carbide structures, we employed the software package JEMS [3], the JAVA version of EMS [4] by Stadelmann, as well as the software package EMS Online [5].
3. Results Fig. 1 presents a conventional bright-field TEM image of a cross-sectional specimen of the foil ADT15. The micrograph reveals heavily dislocated austenite (dark regions), labeled c, intermixed with needles (or laths) of a second phase, labeled v, about 50 nm in diameter and up to several micrometers long. As shown in the fol-
1
Vishay micro-measurements.
Fig. 1. Cross-sectional TEM bright-field image of carbide particles in the carburized foil ADT15.
lowing, these particles are v-carbide. Often, the needles lie perpendicular to the surface. As seen in Fig. 1, they can be present with a higher volume fraction than the austenite. The straight vertical features normal to the needle axes in Fig. 1 indicate the presence of a high concentration of planar faults in the needles. TEM images of KMD-A, KJL-A, and ADT40 reveal similar scenarios: heavily dislocated austenite, containing elongated particles of a second phase, invariably containing a high density of stacking faults or microtwins normal to the long axis. According to the results we obtained by XRD (X-ray diffraction) [1], the carbide particles in Fig. 1 should have the structure of Fe5 C2 (H€agg carbide or v-phase, Table 3), shown in Fig. 2. Attempts to study the crystal structure of the carbide particles by convergent-beam electron diffraction (CBED) remained unsuccessful because the patterns did not reveal any HOLZ lines, or
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Table 3 Crystal structures Carbide
Lattice parameters
Space group
Atom positions
v-Fe5 C2 (H€ agg carbide)
a ¼ 1:1552 nm, b ¼ 0:4564 nm, c ¼ 0:5043 nm, a ¼ c ¼ 90° b ¼ 97:68° obtained by fitting experimental data with Celref software [24]
C2/c (15) monoclinic
Fe in 0.097, 0.078, 0.423; Fe in 0.215, 0.581, 0.306; Fe in 0, 0.561, 0.25; C in 0.107, 0.285, 0.149 obtained from [25]
Fe7 C3 a
a ¼ 0:4537 nm, b ¼ 0:6892 nm, c ¼ 1:1913 nm a ¼ b ¼ c ¼ 90° (JCPDS file 75-1499)
Pnma (62) orthorhombic
Fe in 0.25, 0.07, 0.02; Fe in 0, 0.07, 0.81; Fe in 0.25, 0.25, 0.20; Fe in 0.25, 0.25, 0.42; Fe in 0, 0.25, 0.63; C in 0.11, 0.03, 0.35; C in 0.38, 0.25, 0.57 obtained from [15]
a
Fe7 C3 is isostructural to Cr7 C3 with a ¼ 0:4532 nm, b ¼ 0:7015 nm, c ¼ 1:2153 nm (JCPDS file 36-1482).
Fig. 2. Crystallographic model of Fe5 C2 , the ‘‘H€agg’’ carbide (four unit cells in ½0 1 0 projection).
other three-dimensional information, and the pattern in the discs was too irregular for symmetry considerations – apparently as a result of the high defect density and internal stresses. However, Fig. 3(a), the selected-area electron diffraction pattern of the region shown in Fig. 1, confirms the hypothesis that the carbide particles in Fig. 1 have the structure of Fe5 C2 : According to the simulated diffraction pattern in Fig. 3(b), Fig. 3(a) corresponds to the pattern expected for H€ agg carbide in ½0 1 0 projection. The fine streaks at each spot in Fig. 3(a) apparently originate from the planar faults (stacking faults or twins) normal to the long axes of the carbide needles; our rotation calibration for the CM20 indicates that they are oriented orthogonal to the fault plane. The orientation of the streaks in Fig. 3(a) indicates that the planar faults are parallel to the plane spanned by the unit cell vectors b and c of the monoclinic crystal structure. Accordingly, the long axis of the needles corresponds to the a axis of the (monoclinic) unit cell. The diffraction pattern of Fig. 3(a) also reveals faint reflections from the austenite matrix. According to these reflections, the austenite had a h2 1 1i axis parallel to the incident electron beam, and the following orientation relationship exists between the carbide (index v) and the austenite (index c): h0 1 0iv k h2 1 1ic ; f0 0 1gv k f1 1 1gc :
ð1Þ
The selected-area diffraction pattern of Fig. 3(c), recorded in a different TEM specimen (a cross-sectional specimen of ADT15), reveals the same common diffraction pattern of carbide and austenite as Fig. 3(a). This indicates that the orientation relationship (1) is significant. In fact, this orientation relationship seems to be unique; evaluating the diffraction patterns of about 20 further H€agg carbide particles, we have not observed any other orientation relationship. The particles of Fig. 1 are all within the same austenite grain, and they are all aligned with respect to each other. This means that not only their lattice has a preferred orientation relationship with the austenite lattice, but also that their long axis corresponds to a particular crystallographic direction in the carbide (b c) and in the austenite. Elongated carbide particles, similar to those in Fig. 1, were also observed in plan-view and cross-sectional specimens of carburized KMD-A, KJL-A, and ADT40, the thicker one of the two foil samples. As an example, Fig. 4(a) shows a plan-view dark-field TEM image of carbide particles in a cross-sectional specimen of KMDA. The corresponding selected-area electron diffraction pattern in Fig. 4(b) is consistent with the hypothesis that these particles have the same crystal structure as those in Fig. 1, thus the structure of Fe5 C2 . The pattern corresponds to that of Fe5 C2 in ½0 0 1 projection, and the reflection that served for recording the image Fig. 4(a) is encircled and identified as a ð5 3 0Þ reflection. Not only the structure, but also the morphology of the carbide particles, which show up bright in the dark-field image of Fig. 4(a), closely corresponds to those in ADT15 (Fig. 1). Evaluation of about 50 electron diffraction patterns of needle-shaped carbide particles in all bulk and foil samples agree with the hypothesis that most of these particles have the crystal structure of H€agg carbide. Diffraction patterns consistent with the structure of Fe5 C2 are also observed in TEM diffraction patterns from the powder specimen P42860. According to the results of our XRD studies, reported in [1], the powder is fully converted to carbide. Our TEM studies have confirmed this finding. Fig. 5 shows two examples of small (electron-transparent) powder particles and the corresponding selected-area diffraction patterns. These
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Fig. 4. Carbide particles in KMD-A. (a) Plan-view dark-field TEM image. (b) Corresponding selected-area electron diffraction pattern. The ð5 3 0Þ reflection that served for recording the image (a) is encircled.
Fig. 3. Selected-area diffraction patterns. (a) Pattern recorded in the region shown in Fig. 1, rotated relative to Fig. 1 according to our calibration of image rotation. (b) Simulated diffraction pattern of H€ agg carbide in ½0 1 0 projection. (c) Pattern recorded in a different specimen, indicating the same orientation relationship between the H€ agg carbide and the austenite (c) as in (a).
patterns, recorded in ½1 0 0v and ½1 1 0v direction, exclusively exhibit carbide reflections. The powder particles possess pronounced facets, corresponding to lowindexed crystallographic planes, and the entire morphology – particularly the special angles the facets make with each other – reflects the low symmetry of the monoclinic crystal structure.
In summary, our observations lead to the conclusion that carbide particles with the crystal structure of Fe5 C2 (H€agg carbide) have formed in all five carburized specimens. Since 316-type steels also contain major amounts of Cr and Ni, it seems likely that these elements are also present in the carbide phase. Indeed, high-resolution analytical TEM has confirmed this hypothesis [6]. Based on the results from XRD, the observation of H€agg carbide by TEM was expected. However, we also observed particles of another carbide, which had not been detected by XRD before. The electron diffraction patterns of this carbide are consistent with the crystal structure of Fe7 C3 (Table 3, Fig. 6). Fig. 7 presents a selected-area electron diffraction (SAD) pattern of such a carbide particle in a plan-view, back-thinned specimen of the foil sample ADT40. This pattern exhibits sixfold symmetry, indicating the existence of either a threefold or a sixfold symmetry axis in the crystal structure. Since other diffraction patterns we have obtained from these carbide particles, including the pattern of Fig. 9 discussed below, are incompatible with cubic symmetry, Fig. 7 suggests that this carbide has either hexagonal or trigonal/rhombohedral symmetry. In both cases, the
F. Ernst et al. / Acta Materialia 52 (2004) 1469–1477
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Fig. 7. Selected-area diffraction pattern of a carbide particle in the foil sample ADT40. The inset at the lower left corner shows an enlargement of the area in the small dashed square above. The circled reflections originate from a neighboring austenite grain.
Fig. 5. TEM bright-field images of two powder particles from the carburized powder sample P42860 and corresponding selected-area diffraction patterns. The labels in the diffraction patterns refer to the encircled reflections.
Fig. 6. Crystallographic model of Fe7 C3 (four unit cells in [1 0 0] projection).
structure can be described by a hexagonal Bravais lattice with lattice parameters aC and cC and the viewing direction of Fig. 7 corresponds to the ½0 0 0 1 or ½0 0 0 1 direction of the hexagonal crystal lattice. Fig. 8 presents an example for the particle morphology of second carbide, observed in a cross-sectional TEM specimen of KJL-A. The appearance of these particles closely resembles that of the M5 C2 carbide particles of Fig. 1: They are elongated, needle-shaped, with a diameter of a few tens of a nanometer, and
Fig. 8. Cross-sectional TEM bright-field image of carbide particles the carburized foil ADT15. These particles have a different crystal structure than those of Fig. 1.
contain a high concentration of planar faults (particularly obvious in the arrowed particle, accidentally oriented with the fault planes precisely parallel to the viewing direction). Correlating TEM images and diffraction information (using the image rotation calibration of our instrument) revealed that the ½0 0 0 1 direction of the hexagonal axis coincides with the long axis of the carbide grains of Fig. 8. The six encircled, weak diffraction spots in Fig. 7 are f2 2 0g reflections of adjacent austenite, which was also included in the SAD aperture. Accordingly, the main symmetry axis of the carbide lies parallel to a h1 1 1i
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direction of the austenite. The h1 1 0i directions of the austenite, however, do not coincide with prominent, low-indexed directions in the carbide. As can be seen in Fig. 7 – most clearly in the enlarged inset at the bottom left – each f2 2 0g austenite reflection nearly coincides with a f5 4 1 0g reflection of the carbide (a reflection of the type 5aC þ bC in Fig. 7). It appears that the 3% misfit between the corresponding lattice planes in the carbide and in austenite prevents the carbide from aligning its lattice with the more densely packed f2 2 0g planes of austenite. Consequently, the translations aC and bC in the carbide pattern of Fig. 7 approximately correspond to h5 4 1i directions of the austenite, not h1 1 0i. By evaluating the carbide f1 1 0 0g spacing, corresponding to aC in Fig. 7, relative to the f2 2 0g spacing of the austenite, and by measuring the expanded lattice parameter at the surface of ADT40 by XRD, which yielded ac ¼ 0:3661 nm, we obtained kaC k
1
¼ ð0:616 0:005Þ nm:
ð2Þ
For the lattice parameter a of the carbide phase parallel to the basal-plane, this implies 2 1 aC pffiffiffi kaC k ¼ ð0:711 0:006Þ nm: 3
ð3Þ
(The error limits given here and in the following exclusively reflect the error in determining the positions of the diffraction spot intensity maxima.) Fig. 9 shows another diffraction pattern obtained from a carbide particle of the second kind. This pattern shows pronounced streaks, reflecting the high density of planar faults observed in Fig. 8, and exhibits mm symmetry. Like Fig. 7, the pattern includes diffraction spots from the carbide phase as well as from the neighboring austenite. Since the pattern includes f1 1 1g as well as
Fig. 9. Selected-area electron diffraction pattern of an M7 C3 carbide particle in h2 1 1 0i projection. The bold reflections originate from a neighboring austenite grain in h3 2 1i projection.
f3 3 1g reflections of austenite, the viewing direction corresponds to h3 2 1i in the austenite. The lattice planes producing the fundamental diffraction vector aC lie parallel to f5 4 1g planes in austenite. Measuring the length of aC in Fig. 9 relative to the distances of the austenite spots from 000 and calibrating the austenite spacings with the lattice parameter ac ¼ 0:3661 nm we obtained by XRD for the expanded austenite at the surface of ADT40, we obtained kaC k
1
¼ ð0:613 0:002Þ nm:
ð4Þ
This result deviates by less than 0.5% from the result (2), and therefore identifies the reflecting planes as f1 1 0 0g planes of the hexagonal lattice. The lattice planes producing the fundamental diffraction vector cC lie parallel to f1 1 1g planes in the austenite. According to the information obtained from Fig. 7 and in agreement with the 90° angle they make with the f1 1 0 0g planes, the planes causing the reflection cC constitute the basal planes of the carbide. Hence, the direction in which the pattern of Fig. 9 was obtained corresponds to a ½2 1 1 0 direction of the hexagonal lattice. For the plane spacing that corresponds to cC we obtained kcC k
1
¼ ð0:455 0:002Þ nm:
ð5Þ
At this point, we can justify our assumption that the structure of the second kind of carbide is either hexagonal or rhombohedral, but not cubic. Consider the ratio of the plane spacings (4) and (5) q :¼
kcC k1 kaC k
1
¼ ð0:739 0:004Þ
ð6Þ
If the crystal structure was cubic, Fig. 7 would be the diffraction pattern for a h1 1 1i direction of the primary electron beam, and the fundamental reflections (corresponding to aC ) would be of the type fh h 0g. The pattern of Fig. 9 would correspond to a primary beam direction orthogonal to one of these reflections, for example orthogonal to ½h h 0. Then, the reflections corresponding to cC must have the form ðk k lÞ. Usually, h ¼ 1 for SC (single cubic) or BCC (body-centered cubic) structures and h ¼ 2 for a FCC (face-centered cubic) structures. However, there are no sets of planes in SC/BCC and FCC crystals that would yield a q fulfilling the condition (6). The closest match is obtained for ðk k lÞ ¼ f0 0 2g in SC/BCC and ðk k lÞ ¼ f0 0 4g in FCC, yielding q ¼ 0:707 in each case, but the discrepancy with our result (6) clearly exceeds the error limits. It may be argued, of course, that the 1 1 0 reflections (in SC/BCC) or 2 2 0 reflections (in FCC) could be kinematically forbidden. Under this assumption, it is indeed possible to find plane spacings that match with our result (6). For example, if the fundamental reflections in Fig. 7 were not f2 2 0g but f4 4 0g reflections of a hypothetical FCC structure, the reflections corresponding to cC could be of
F. Ernst et al. / Acta Materialia 52 (2004) 1469–1477
the type f5 5 3g, yielding q ¼ 0:736. However, this would imply not only a primary beam direction with very high indices, h3 3 1 0i in this case, but also a very large lattice parameter: 3.47 nm according to (4) and (5). Considering that Cr23 C6 , for example, has a lattice parameter of 1.07 nm and 116 atoms in the unit cell, this would mean that there are about 4000 atoms in the unit cell. While not impossible, this is at least highly unlikely. Assuming further extinctions in the h1 1 1i zone lead to even more unrealistic results. Therefore, we have discarded the possibility that the observed carbide of the second type has a cubic structure. The diffraction pattern of Fig. 10, obtained from another cross-sectional TEM specimen of ADT40, confirms the lattice parameters and the special orientation relationship deduced from the patterns in Figs. 7 and 9. Summarizing, the crystallites of the second carbide phase also adopt a special orientation relationship with the residual austenite h0 0 0 1iC k h1 1 1ic ; h1 1 0 0i k h5 4 1i ; C
ð7Þ
c
which implies h2 1 1 0iC k h3 2 1ic :
ð8Þ
While the close-packed planes of both phases are parallel, the close-packed directions are not, and there are no common mirror planes of the two adjacent crystal structures parallel to the h0 0 0 1iC k h1 1 1ic direction. According to the diffraction pattern in Fig. 7, the orientation relationship (7) optimizes the alignment of f5 4 1 0gC planes with f2 2 0gc planes, which have nearly the same spacing.
Fig. 10. Selected-area electron diffraction pattern of an M7 C3 carbide particle in h2 1 1 0i projection, similar to Fig. 9, but recorded in a different specimen. Like in Fig. 9, the bold reflections originate from a neighboring austenite grain in h3 2 1i projection, confirming the significance of the corresponding orientation relationship.
1475
Among the relevant iron carbides, the lattice parameters (2) and (5) only match well (within the limits of error) with those of Fe7 C3 as given by Herbstein and Snyman [7–9]. While the XRD data of those authors suggested a hexagonal structure with a ¼ 0:6882 nm;
c ¼ 0:454 nm;
ð9Þ
later work of Fruchart and Rouaoult [7] revealed that this carbide actually has orthorhombic crystal symmetry (Fig. 6) with the lattice parameters a ¼ 0:454 nm;
b ¼ 0:6879 nm;
c ¼ 1:1942 nm;
ð10Þ pffiffiffi implying c=b 3. This structure was confirmed by TEM studies of Morniroli et al. [10–13]. These authors found that M7 C3 carbides, including Fe7 C3 , Cr7 C3 , and Mn7 C3 , tend to contain a high density of planar defects, mainly twin and antiphase boundaries, between the three different orientation variants of the orthorhombic unit cell. It is because of these defects that the material appears to have hexagonal crystal symmetry under the typical experimental conditions of electron diffraction in a transmission electron microscope. A high density of planar defects seems to be an intrinsic feature of Fe7 C3 and was in fact already reported by Dyson and Andrews in 1969 [14]. With reference to the true, orthorhombic crystal lattice, the viewing direction of Fig. 7 corresponds to ½1 0 0 and the viewing direction of Figs. 9 and 10 to ½0 1 0. While the lattice parameters of Fe7 C3 , Cr7 C3 , and Mn7 C3 [15] are too similar to be safely distinguished under our experimental conditions. Fig. 8 demonstrates that locally the volume fraction of the M7 C3 carbide can be as high as 75%. It seems unlikely, therefore, that this carbide is devoid of iron. On the other hand, analytical TEM has revealed that the carbide particles are not pure Fe7 C3 , either. As will be detailed elsewhere [6], X-ray
Fig. 11. Selected-area electron diffraction pattern of carbide particles in KJL-A.
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energy-dispersive spectroscopy indicates that the major cation species is iron, but the carbides also include significant concentrations of the substitutional alloying elements (Cr, Ni). Another example for the general formation of carbide particles consistent with the Fe7 C3 structure is shown in Fig. 11, which presents a selected-area diffraction pattern of carbide particles in KJL-A. This diffraction pattern corresponds to the structure of Fe7 C3 viewed in ½2 1 1 direction.
4. Discussion To our knowledge, carbide formation in low-temperature-carburized austenitic steels has not been observed before, and the formation of H€ agg carbide (M5 C2 ) in austenitic stainless steel is very unusual. A substantial database exists, on the other hand, about the formation of M5 C2 and M7 C3 carbide under other experimental conditions. In ferritic steels, Dirand and Alqier [16] have observed Fe5 C2 as part of a transformation sequence that leads from martensite to a phase equilibrium of ferrite (a-Fe) and cementite (h-Fe3 C) [16]. In this sequence, v-Fe5 C2 forms from g-Fe2 C and subsequently decomposes into h-Fe3 C. The transformation v ! h occurs progressively through variations in the periodicity of twinning, while the slight change in stoichiometry is apparently accommodated by incorporating planar defects. K€ oninger et al. [17] and Hammerl et al. [18] have observed exactly the opposite sequence of carbides on tempering carbon-rich Fe–C solid solutions prepared by ion implantation. The reason for the observed transformation sequences and the topotaxial orientation relationships (‘‘lattice-invariant deformation’’) between subsequent carbides is seen in the structural similarity between the carbide phases g, v, and h: the carbon atoms reside in triangular prisms of iron atoms, and the individual carbide structures just slightly differ by the patterns in which these prisms are arranged. Given the structural similarity between g; v and h carbide, we speculate that annealing of the low-temperature-carburized material will readily lead to the formation of h (M3 C). On the other hand, the observations reported in the literature raise the question whether the M5 C2 carbide we observe perhaps forms via a precursor of g carbide (M2 C). From our experimental data, however, there is no evidence for the formation of g carbide in low-temperature carburized austenitic stainless steels. M7 C3 is unique in that it has the highest carbon content among the carbides considered here. Also this carbide is built up from structural units of carbon atoms near the center of trigonal prisms of Fe atoms [19]. At present, it is not clear whether what controls the ratio of
M7 C3 to M5 C2 in our material. It seems possible that M5 C2 actually transforms to M7 C3 , not M3 C, under the far-from-equilibrium conditions of low-temperature carburization. In this case, the volume fraction of M7 C3 , which is so much smaller than that of M5 C2 , should increase with increasing time. Tempering experiments are being carried out in our group to further investigate this point. The observation that planar faults in the M5 C2 carbide occur on the planes spanned by the unit cell vectors b and c agrees well with the building principle of the structure discussed by Jack and Wild [19], although we believe that the defect structure reported in the present publication has not been observed before. The formation of these faults may be related to transformation stresses or to positioning errors on the growth face. In M7 C3 carbides, however, a high density of planar defects has also been observed in single crystals – absent transformation-related stresses [10,12,20]. For the carbide particles of both phases, M5 C2 and M7 C3 , preferred orientation relationships exist with the austenite, and the long axis of both types of carbide particles is aligned with a particular crystallographic direction. Crystallographic orientation relationships have often been interpreted in terms of minimization of lattice mismatch between the adjacent phases. Whether this interpretation is correct for the extreme non-equilibrium conditions of carbide formation in austenite with a colossal supersaturation of carbon will be subject of future studies, for example by analyzing the observed orientation relationships with the model of Pirouz et al. [21–23]. Another reason for the establishment of a unique orientation relationship could be the atomistic mechanism of carbide nucleation, as discussed for Fe7 C3 by Dyson and Andrews [14], for example a diffusion-less nucleation mechanism involving shear. Concerning the properties of low-temperature-carburized austenitic stainless steel, the formation of carbides is generally undesirable. While carbides are known to increase the flow stress of austenitic stainless steels, our previous results [1] demonstrate that precipitation hardening contributes only marginally to the surface hardness of low-temperature-carburized 316-type austenitic stainless steels. Moreover, since the carbide particles have different mechanical and electrochemical properties than the austenite matrix, their presence may compromise the fatigue and corrosion behavior. The micromechanism of carbide formation is a subject of ongoing research in our group.
5. Conclusions Our TEM studies austenitic stainless steels carburized at low temperature have provided important insight in the microstructural changes that accompany the carbu-
F. Ernst et al. / Acta Materialia 52 (2004) 1469–1477
rization process after prolonged treatment. In addition to the colossal supersaturation of the austenite with carbon and the formation of an M5 C2 carbide, which we have already observed by XRD [1], TEM has revealed a small fraction of a second carbide phase, M7 C3 , and a high dislocation density in the austenite matrix. We conclude that a quantitative understanding of the surface properties resulting from low-temperature carburization must take into account the entirety of all these features. The formation of carbides under the far-fromequilibrium conditions of colossal supersaturation with carbon constitutes a very unusual scenario, which warrants further analysis.
[4] [5] [6] [7] [8] [9] [10]
Acknowledgements
[18]
We thank Peter Williams, Sunniva Collins, and Steven Marx for numerous fruitful discussions, Arthur Heuer and Fumiyasu Oba for comments on the manuscript, and the Swagelok Company for financial support.
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