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Deformation-induced microstructure and martensite effects on transgranular carbide precipitation in type 304 stainless steels
Acta metall, mater. Vol. 41, No. 9, pp. 2589-2600, 1993
0956-7151/93 $6.00.4-0.00 Copyright © 1993Pergamon Press Ltd
Printed in Great Britain. All rights reserved
DEFORMATION-INDUCED MICROSTRUCTURE A N D MARTENSITE EFFECTS ON TRANSGRANULAR CARBIDE PRECIPITATION IN TYPE 304 STAINLESS STEELS A. H. ADVANI, L. E. MURR, D. J. MATLOCK, R. J. ROMERO, W. W. FISHER, P. M. TARIN, J. G. MALDONADO, C. M. CEDILLO, R. L. MILLER and E. A. TRILLO Department of Metallurgical and Materials Engineering and Materials Research Institute, University of Texas at El Paso, E1 Paso, TX 79968, U.S.A. (Received 18 September 1992; in revised form 5 March 1993)
Abstract--Plastic deformation of 304 stainless steel (SS) induces transgranular (TG) carbide precipitation, which is critically dependent on deformation-induced microstructural changes occurring during thermal treatment of the SS. Uniaxial deformation of the 304 SS to 40% strain produces a high density of intersecting micro-shear bands composed of heterogeneous bundles of twin-faults and about 12-17% strain-induced ct'-martensite at the intersections of the twin-faults. Thermal treatment at 670°C for 0.1-10 h, however, results in a rapid annihilation/transformation of the strain-induced martensite and the concurrent formation of zones containing mixed thermal martensite laths and fine-grained austenite, though the thermal martensite also decreases with increasing heat treatment time. Simultaneous with these thermomechanically-induced microstructural changes, TG chromium-rich carbides form at intersections of twin-faults and on fine-austenite or thermal martensite boundaries in the SS; however, no correlation between strain-induced ~'-martensite and carbides was observed in this work. The mechanisms of deformation-induced microstructure and (strain-induced and thermal) martensite effects on TG carbide precipitation in 304 SS are discussed.
INTRODUCTION Sensitization refers to the development of grain boundary chromium-depletion caused by the precipitation of chromium-rich carbides on grain interfaces of stainless steel (SS) [1]. Sensitization occurs during thermal or thermomechanical (TM) treatment of SS, especially when the material is isothermally aged or continuously cooled through the temperature range of 500-850°C. Chromium-rich carbides precipitate on the grain boundary during these TM treatments, and due to a relatively slow diffusivity of chromium vs carbon, a substantial amount of chromium is drawn from the region around the growing carbide. This leads to grain boundary chromium-depletion, or sensitization, in the SS and renders the material susceptible to intergranular (IG) corrosion and IG stress-corrosion (SCC) cracking in certain aqueous environments. Example environments where SS is susceptible to IG corrosion and IGSCC include high temperature oxygenated water used in boiling water reactors, or corrosive environments present in containers used to store and transport toxic and radioactive waste materials where the temperatures can rise above ambient [2, 3]. Plastic deformation accelerates the development of grain boundary carbide precipitation and sensitization in 304 and 316 SS [3-17]. In recent work on 316 SS, we have shown that the acceleration of sensitization is a function of the amount of (prior) strain in the material and temperature of isothermal sensitiz-
ation treatment [3-5]. A systematic increase in strain from 0 to 20% was observed to yield a continuous increase in sensitization below 700°C, while at higher temperatures, a threshold strain of 5-10% was required to produce accelerated chromium depletion. Strain-induced acceleration in sensitization was also indicated to be caused by an acceleration in the kinetics of chromium diffusion, and by a lowering of the free energy barrier to carbide nucleation in the SS [3, 5]. Straining above 20% produces higher amounts of chromium depletion through the simultaneous precipitation of IG and transgranular (TG) (or grain matrix) carbides in 304 and 316 SS [12-19]. Transgranular chromium-depletion in 316 SS has been observed to be strain and heat treatment dependent, yielding C-curve behavior on a timetemperature-strain-TG corrosion map [18, 19]. The strain effect dominates, and increased amounts of TG depletion, and decreased time to develop TG corrosion occur with increasing strain level. Higher temperatures also reduce the time to develop grain matrix attack, while isothermal holding for longer times produces higher amounts of T G corrosion in SS. Transgranular depletion in 316 SS has been shown to occur due to precipitation of carbides within grain-matrix regions [18, 19]. Specifically, defect sites created during straining are favored locations for T G carbide precipitation in SS [19-31]. Transmission electron microscopy (TEM) has revealed that regions of high dislocation density,
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ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
deformation twins, stacking-faults, and twin-fault intersections are preferred sites for matrix precipitation in deformed SS. Briant and Ritter [7-9] have compared the effects of strain on sensitization in 304 and 316 SS deformed to 5 ksi below the failure stress of the material, and indicated fundamental differences in deformation-induced sensitization for these materials. Briant and Ritter showed that 304 SS had accelerated T G and IG corrosion over 316 SS, and the amount of T G attack was especially more pronounced in 304 vs 316 SS. This was attributed to rapid precipitation of carbides, and development of chromium-depletion, on straininduced martensite and martensite lath boundaries in 304 SS, vs on slip planes in 316 SS. However, the evidence provided to relate the martensite to carbide precipitation only showed carbides on martensite laths, and the relationship between strain-induced martensite and carbide precipitation was not documented. To our knowledge, the evolution of deformation microstructure, martensite transformation, and T G carbide precipitation with TM treatment has also not been defined in any work to-date. Since the mechanisms and crystallography of strain-induced martensite formation in 304 SS are known [32-37], we believed that an understanding of TM processing effects on deformation microstructure (especially martensite), and identification of key relationships between strain-induced martensite and T G carbide precipitation, would be critical in evaluating mechanisms of carbide precipitation and chromiumdepletion phenomena in SS. The objective of this work was to develop a fundamental understanding of deformation-induced microstructure and martensite effects on T G carbide precipitation in 304 SS. The work was designed to focus on using TEM to investigate the relationship between T G carbides, deformation-induced sites and strain-induced martensite in 304 SS. Bulk X-ray diffraction (XRD) and magnetic measurement techniques were used to support the observed changes in microstructure during TM treatment of 304 SS. EXPERIMENTAL TECHNIQUES Material, deformation processing and heat treatment The material used in this work was a high carbon, Type 304 austenitic stainless steel of the composition shown in Table 1. The material was received as a 6.25 mm thick plate in the mill processed condition, and solution annealed at I100°C for 1 h prior to deformation processing. The solution annealed material had large, equiaxed grains, with an average grain-size of 175 gm. The microstructure of the annealed material was also
%C 0.051
confirmed to be clean, with a relatively low-dislocation density and no carbide precipitation on grain boundaries, which made it an appropriate initial condition for the deformation and heat treatment experiments. The solution annealed plate was cut into strips of length 125mm, and width 12.5mm. The sample was marked at 12.5mm intervals, and the initial a r e a (Ai) was measured at the marked locations on the sample. The strips were pulled in uniaxial tension using an Instron universal testing machine. The final area at the marked locations on the sample was then measured (Af), and the true strain was calculated using the relation true strain = In (Ai/Af). In this work, samples were deformed to ~ 4 0 % true strain, which corresponds to an engineering strain of ~49%. The deformed samples were cut into sections of 12.5 x 12.5 mm, and then heat treated at 670°C for 0.1-10 h to examine the effects of strain and heat treatment on carbide precipitation in SS.
Examination o f deformation microstructure and carbide precipitation in S S The deformation microstructure and carbide precipitation in 40% true strain, 670°C/0-10 h samples of SS were examined using a Hitachi H-8000 scanning transmission electron microscope (STEM), operated at 200 kV. Conventional bright-field TEM, dark-field TEM, selected-area electron diffraction and microdiffraction techniques were used for identification of phases produced during the TM treatments. Limited energy dispersive X-ray spectroscopy (EDS) was also used to identify chromium-rich carbide phases in SS. The formation of martensite in deformed samples, and the effect of heat treatment on the martensite transformation were also examined using XRD and magnetic measurement techniques. The austenite-to-martensite transformation was qualitatively studied using X R D by examining and comparing the face-centered cubic (f.c.c.) austenite peaks vs the body-centered cubic (b.c.c.) martensite peaks in the X R D spectrum. Also, since the martensite product is magnetic, the austenite-to-martensite transformation is ferromagnetic and was readily detected by magnetic measurements. The magnetic permeability of 40% strain, 670°C, 0-10 h samples was measured with a commercially available Ferritescope, which is calibrated to measure the content of ferrite in weldments. Additional details of the magnetic measurement technique to determine% martensite in SS are described elsewhere [37].
Table 1. Chemicalcompositionof type 304 stainlesssteel heat (in wt%) %Mn %P %S %Si %Ni %Cr %Mo %Co %Cu 1.47 0.025 0.014 0.38 8.22 18.29 0.19 0.11 0.19
%N 0.085
ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL RESULTS Plastic deformation effects on microstructural development
Plastic deformation to 40% strain produced several microstructural features in 304 SS, the most significant of which are illustrated in Fig. 1. The figure shows a bright-field TEM micrograph, a dark-field TEM micrograph and a microdiffraction pattern which illustrate the formation of strain-induced ~'martensite at twin-fault intersections in the as-deformed SS. The bright-field micrograph indicates the presence of intersecting microshear bands in the deformed microstructure of the 304 SS. These microshear bands are usually composed of heterogeneous bundles of stacking faults and deformation twins, and in some cases may also include E-martensite. Straininduced martensite is seen to form within the twinfault intersections; locations of the martensite are clearly indicated by the arrows in the bright field micrograph, and also by the "lit-up" areas in the dark-field micrograph. The presence of strain-induced martensite was confirmed by microdiffraction analysis of the twin-fault intersections which contained martensite.
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BCC c~' (I11) Fig. 1. Illustration of strain-induced martensite formation in 304 SS. The TEM micrographs are bright-field and darkfield views of the strain-induced ~'-martensite at twin-fault intersections in the SS. The arrows in the bright-field view show some locations of the ~'-martensite in the micrograph. The diffraction pattern results by microdiffraction from a martensitic intersection. AMM 41 t~-F
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Only a fraction of the twin-fault intersections in the SS was observed to transform to strain-induced martensite. Bulk ferromagnetic analysis, in fact, quantified the fraction of ~'-martensite in the 40% strain, as-deformed 304 SS to be 12-17% (depending on the orientation of measurement). X-ray diffraction analysis also qualitatively confirmed that a fraction of the 304 SS material was transformed to strain-induced martensite, as can be seen in a comparison of relative peak heights for ~'-martensite and y-austenite peaks shown in Fig. 2. Thermal treatment effects on defi~rmation microstructure
Heat treatment of the 40% true strain 304 SS produced significant and rapid changes in the deformation microstructure of the material. For the shortest time of heat treatment examined, i.e. 0.1h at 670°C, no strain-induced martensite was observed at twin-fault intersections, unlike that seen in the asstrained material. Instead, a "new" cell-like phase was seen to form at the twin-fault intersections, as is shown in the bright-field micrographs and microdiffraction patterns of Fig. 3(a) and (b). The high magnification view of Fig. 3(a) illustrates the elongated cellular morphology of the phase. Analysis of the microdiffraction of Fig. 3(b) indicated that the phase was b.c.c.-martensite. However, the cell-like appearance of the phase, and the extent of its growth even outside twin-fault intersections suggested that the phase was not strain-induced martensite, but was more typical of lath martensite; hence we considered the phase to be "newly" formed during the heat treatment. It was also apparent that the lath martensite formed during the heat treatment process, while the strain-induced martensite was seen only in (room temperature) strained SS; hence, we considered the lath martensite to be a "thermally" induced, rather than a strain-induced, product. And, to assist in our differentiation between the two martensites, we have chosen to use the terms "lath martensite" and "thermal martensite" to refer to the thermally-induced phase, while the "strain-induced-martensite" and "~'-martensite" will continue to imply the strain-induced phase. Other locations which appeared to be formed by a continued production of the thermal martensite were also observed in this 670-C-0.1 h sample (Fig. 4). The cell-like regions again originated within the twin-fault intersection regions, as is illustrated in the bright-field and dark-field micrographs of Fig. 4. Selected-area electron diffraction indicated a complex ring pattern, which is possibly due to the fine-intermix of thermal martensite dispersed within the austenite at the location. Longer heat treatments to 0.4 h at 670'~C resulted in an increased observation of the clustered bands of cellular phases, and in several locations the clusters often encompassed entire sections of grains (Fig. 5). The clustered cells appeared to be oriented in specific
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ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 2. X-ray diffraction spectrum showing martensite and austenite peaks in the spectrum. The ~t' represents the strain-induced martensitic peaks, while the ? represents the austenitic peaks. directions, and often appeared as bands in the material. Although some clusters were also observed in the 40% strain, 670°C-0.1 h sample, the extent of clustering was more prominent in the 0.4 h sample. Selected-area electron diffraction again indicated the presence of a complex ring-pattern in the region. Further microdiffraction analysis was, thus, essential within the clusters in this sample to confirm the phases within the cell-locations. Microdiffraction analysis of the clustered region indicated that the cellular region was not entirely thermal martensite, but instead contained a mixture of thermal martensite and fine-grained austenite (Fig. 6). We will refer to this fine-grained austenite that forms within local, T G clustered regions as "fine-austenite" to distinguish it from the largegrained austenite in the initial microstructure. An example dimension of the fine-austenite grain-size was measured to be 0.2 x 0.3 #m, in comparison with the original solution-annealed grain size of 175 #m. Both the fine-austenite and lath martensite are seen to have relatively low-dislocation densities in Fig. 6. The fine-austenite may also be contributing to the complexity of ring selected-area electron diffraction patterns observed in Figs 4 and 5. Additional heat treatment to 1 and 10 h at 670°C continued to show the presence of clustered cellular regions in the microstructure. No significant changes in the fine-austenite size or lath martensite size were seen from the 0.1-I0 h of treatment. Microdiffraction
of several locations again indicated that both fineaustenite and martensite were present within the clusters, although it was considered possible that there may be changes in the amount of martensite vs austenite ratios in these samples. To verify this possibility, we used bulk magnetic measurements and X R D analysis of all heat treated samples. Ferromagnetic analysis indicated that the amount of martensite changed dramatically with longer heat treatments at 670°C. The martensite was relatively constant (about 12-17%) from 0.0 to 0.4 h of heat treatment, and then dropped sharply to 2-3% after 1 and 10 h at 670°C (Fig. 7). The detection of the martensite by magnetic techniques was also seen to have an orientation effect, which is not understood. XRD analysis supported the bulk magnetic observations and indicated the same qualitative trends in martensite level and an orientation effect of the martensite, though it was difficult to quantify the martensite during X R D (Fig. 8). Thus, both these bulk techniques indicate that the clustered "cellular" regions seen in the 1 and 10 h samples are almost always austenite, and our observations of some lath martensite in the TEM samples is probably a statistical effect. Thermal treatment effects on transgranular carbide precipitation
Heat treatment also resulted in T G carbide precipitation, which occurred simultaneously with the
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 3. Formation of thermal martensite laths at twin-fault intersections in the 40% -670°C4). 1 h sample.
Fig. 4. Formation of clustered regions at twin-fault intersections in the 40%-670°C~0.4h sample. Micrographs represent bright-field and dark-field views of the clusters. The selected-area electron diffraction pattern shows a complex ring diffraction pattern is obtained from the location.
Fig. 5. Illustration of the enhancement of the cluster~l region in the 40% strained sample with 0.4 h of treatment at 670°C. The clusters are seen to have a specific orientation, and the diffraction pattern is complex showing a ring pattern (and texturing effects) due to the fine intermixing of thermal lath martensite and untransformed austenite.
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CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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deformation-induced microstructural changes in the SS. It should be noted that while carbides were also observed to form on grain boundaries in the 40% deformed SS, we have centered on the understanding of T G carbide precipitation in the SS, which is the focus of this work. Two primary locations for T G carbide precipitation were observed in this work: (1) on the boundaries within the mixed fine-austenite and martensite lath regions; and (2) on intersections of twin faults in the SS. An example view of an isolated carbide at a region of fine-austenite and martensite in the 0.4h sample is shown in Fig. 9. It illustrates the complexity of determining the presence of T G carbides in these materials, and indicates the key role of microdiffraction in the analysis• The brightfield TEM micrograph and corresponding selectedarea electron diffraction pattern did not provide a clear indication of the phases present at the location• However, microdiffraction at locations 1, 2 and 3 clearly reveals the presence of austenite, lath martensite, and carbides at these locations, respectively. It can be seen that the fine-austenite and martensite regions are completely intermixed. The carbide seems to form at a martensite lath boundary/fineaustenite boundary within the region• Additional evidence for the formation of carbides on the boundaries of mixed, fine-austenite grains and thermal martensite laths in these materials is provided in Fig. 10.
In the samples heat treated beyond 1 h and up to 10 h at 670°C (where only some thermal martensite was present), carbides were continued to be observed within the clustered regions. This is illustrated in the TEM micrographs showing lattice images of carbides formed in a predominantly fine-austenitic region (Fig. 11). Detailed views of the lattice images show undistorted carbide lattices, dislocations at the carbide-matrix interface, and changes in lattice orientation across the fine-austenitic interface. The other important location for T G carbides precipitation was on twin-fault intersections in the SS. Transgranular carbides on the intersections were difficult to observe, especially in the short heat-treatment (0. I, 0.4 h) samples• This is probably due to a small size of carbides, low density of carbides, or due to a well-spread distribution of carbides within the grain• Larger carbides were, however, also seen in the longer heat treatment (670°C-10 h) samples• Examples of small carbides on twin-fault intersections in the 0.4 h sample are shown in Fig. 12, while larger carbides seen at twin-fault intersections in the 670°C-10h sample are illustrated in Fig. 13. No strain-induced martensite was seen in the vicinity of the carbides in these or other twin-fault locations of the SS. DISCUSSION
Plastic deformation produces strain-induced ~t'martensite formation at the intersections of twin-faults in as-strained 304 SS. Subsequent heat treatment results in an annihilation/transformation of the strain-induced martensite, and in the simultaneous nucleation of thermal lath martensite, though the lath martensite also eventually decreases with increasing heat treatment time. The strain-induced martensite forms via the specific 20
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Fig. 7. Ferromagnetic analysis indicating variations in the amount of martensite with heat treatment. It should be noted that the amount of martensite measured was dependent on the orientation of the sample (face or edge) on which the martensite was recorded. The low martensite values of 2-3% should also be noted to be inaccurate as they a r e typically below the detection limit of the instrument. (Data courtesy Dr K. P. Staudhammer, Los Alamos National Laboratory.)
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Fig. 9. An example view of an isolated TG carbide at a region of fine-austenite and thermal martensite in the 40%--670°C-0.4 h sample. The bright-field image is on the top-left of the figure and a selected-area electron diffraction taken from the entire region is to the right of the bright-field. Microdiffraction patterns for fine-austenite, thermal martensite and TG carbide (bottom left to right) were taken from locations I, 2 and 3 marked on the bright-field micrograph respectively.
crystallographic conditions defined by Olson and Cohen for 304 SS [32, 33]. The volume of strain-induced martensite formed is also strain-dependent, and the 12-17% transformation observed in this work, reflects the 40% strain in the material [37]. The requirement for a thermal impetus to form the lath martensite has not been previously documented in deformed SS, although some changes in martensite morphology were reported by Murr and Staudhammer in shock-loaded 304 SS [38]. And although the lath martensite and strain-induced martensite phases have the same b.c.c, crystal structure and lattice parameter in SS, we believe that the mechanisms and crystallography of transformation for the two martensites would be significantly different. While no direct evidence of a relationship between strain-induced ~t'-martensite and the lath martensite was noted, the results allude to a possible transformation from strain-induced martensite to lath martensite, for the following reasons. First, both phases formed on twin-fault intersections in the SS, and the observation of the lath martensite coincided with the disappearance of the strain-induced martensite, all of which occurred within a short duration of 0.1 hr at 670°C. And second, the bulk observations also sup-
port the possibility of a complete transformation from strain-induced martensite to thermal lath martensite, This is because the total volume of martensite was essentially unchanged between the as-deformed and 670°C-0.1 and 0.4 h samples, and only thermal martensite was seen in the heat treated samples while only strain-induced martensite was seen in the as-deformed samples. However, the evidence is not conclusive, and it is also possible that the strain-induced martensite is annihilated at twin-fault intersections during the heat treatment, while the lath martensite forms at other unrelated twin-fault intersections within the same period. The clustered regions are produced by a fine intermixing of thermal martensite laths with fine-austenite, and may also include untransformed austenite grains. The clusters form within 0.1 h of heat treatment, though the thermal lath martensite eventually decreases after 1 h at 670°C. The formation of the fine-austenite could be caused by a local-recrystallization process, simultaneous transformation of straininduced or lath martensite to fine-austenite, or local encompassment of untransformed austenite by the thermal martensite laths. And, the lack of growth within fine-austenite grains observed with heat
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 10. Formation of TG carbides on the boundaries of mixed, fine-austenite grains and thermal martensite laths in 304 SS.
treatment may be because the 670°C temperature does not provide sufficient thermal impetus for the growth of these regions, or because the grains are pinned by carbides which form on their boundaries. It should be noted that the T G clustered regions discussed above are similar to the lath martensite regions previously described by Briant and Ritter [7-9]. However, in this work, we have provided a clearer definition of the formation of the region, and conclusively demonstrated that it consists of a fine intermix of thermal lath martensite and fine-austenite, rather than only lath martensite (or strain-induced martensite) as was described previously. Transgranular carbide precipitation is critically dependent on the formation of the clustered regions in the SS, which matches similar observations made by Briant and Ritter [7-9]. The formation of T G carbides on boundaries within the mixed fine-austenite and martensite lath regions is a reflector of the high energetics associated with ferrite-austenite interfaces; in fact, Stickler and Vinckier [39] have shown that carbides prefer to precipitate first on ferrite-austenite boundaries, in comparison with grain boundaries, twin-boundaries, and other deformation sites in SS. However, the loss of thermal martensite after 1 h of heat treatment does not seem to influence the T G carbide precipitation, since the precipitation occurs even at the longer heat treatments; the precipitation on the now completely fine-
2597
austenite regions would be similar to carbide formation on grain boundaries in SS. Thus, while the thermal martensite is a contributor to the TG carbide precipitation process, it is not a necessary requirement for the precipitation to occur within the clustered zones in the SS. The observed lattice images of T G carbides within the predominantly fine-austenitic cluster zone observed in the 670°C-1 h sample indicate key mechanisms of carbide precipitation on grain boundaries in SS. The lattice changes associated with the carbide precipitation in our work on 304 SS have also been reported by Young-Hua et al. [40] during precipitation of M23C6 in a superalloy. The undistorted lattice image that was observed is indicative of a possible coherent match-up between the matrix and precipitate, in which the fringes may pass the carbide-austenite interface undisturbed into the austenite matrix; although no lattice fringes were observed in the matrix of our samples, Yong-Hua et al. have observed this result, and attributed it to the lattice spacing of M23C6 being approximately three times the spacing of the matrix. The dislocations seen on the carbide-austenite interface are due to the carbide trying to minimize the free energy of transformation by maintaining a semi-coherent interface with the matrix, in which the disregistery is taken up by misfit dislocations at the interface [41]; the occasional presence of misfit dislocations was also noted in YongHua's work. Furthermore, the change in lattice orientation across the interface has been suggested [40] to be caused by the carbide nucleating separately on either side of the grain boundary, so as to maintain a high degree of coherency with the grain in which it grows. Thus, there are a range of classical precipitation mechanisms by which the carbide grows in 304 SS, and we believe that the preferred mechanism is specific to each precipitate, so as to minimize the free energy of the system. Transgranular carbides also formed on twin-fault intersections in the 304 SS. However, no evidence of any relation between strain-induced martensite and T G carbides was observed in this work, except that they both preferred to form on twin-fault intersections in the SS. This observation is in contradiction with previous results reported by Briant and Ritter [7-9], and the differences could be due to heat-to-heat variations in SS, TM processing differences, or limited area observed during TEM. While there may or may not be a correlation between strain-induced martensite and T G carbides, it is clear that the martensite phase is definitely not a precursor to the formation of carbides on twin-fault locations in SS. This is supported by our observation of carbide precipitates at twin-fault locations where there was no strain-induced martensite. The absence of a requirement for strain-induced ~'-martensite to be a precursor to the formation of T G carbides is further supported by the fact that carbides also form on twin-fault intersections in 316 SS [18, 19], where the
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CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
Fig. 1I. TEM micrograph showing lattice images of TG carbides formed in a fine-austenitic region. Detailed views of the lattice images, in fact, show the undistorted carbide lattice, dislocations at the carbide-matrix interface, and changes in lattice orientation across the fine-austenitic interface.
~53
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Fig. 12. Example views of TG carbides on twin faults and their intersections in the 670°C4).4 h sample (a, b). No strain-induced martensite was seen to be present in the vicinity of the carbide.
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 13. Illustration of carbides seen at TG twin-fault intersections in the 670°C-10 h sample (a) and (b).
strain-induced martensite transformation does not Occur.
Thus, the key mechanisms that define the role of deformation-induced microstructure and martensite formation on TG carbide precipitation have been identified and fundamentally evaluated in this work. It is clear that it is not only plastic deformation, but the synergistic effect of the deformation and thermal treatment, which produce the required microstructural sites that are key to the precipitation of the T G carbides in the SS. The strain-induced martensite phase which lbrms in the as-deformed material is annihilated/transformed during the thermal treatment and, therefore, has no correlation with the carbide precipitation in SS. Instead clustered zones that consist of thermal lath martensite and fineaustenite, and twin-fault intersections (with no martensite), are more potent sites for the T G carbide formation, which occurs by classical nucleation mechanisms to minimize the free-energy of the system. And, while the lath martensite contributes to the TG precipitation process, neither the strain-induced martensite nor lath martensite are a precursor or required for carbide precipitation in the SS. All these evidences point to the fact that the energetics, rather than the crystal structure, of the site is a key to the precipitation of TG carbides in the SS. Ultimately, the energetically driven, deformation site-specific behavior of TG carbides will control the sensitization,
corrosion and SCC behavior of heavily deformed SS in service applications. SUMMARY AND CONCLUSIONS QPlastic deformation of 304 SS to 40% strain produces microshear bands composed of twin-faults and about 12-17% strain-induced martensite. QThermal treatment of the deformed SS results in the annihilation/transformation of the strain-induced martensite, and the simultaneous formation of mixed regions containing about 12-17% thermal martensite laths and fine-austenite. The lath martensite formation also occurs at twin-fault intersections in the SS. Longer heat treatment (beyond 670°C-1 h) results in a reduction in the amount of thermal martensite in the SS to 2-3%. • Transfer anular carbides form on the mixed, thermal martensite lath and fine-austenite boundaries in the SS. They also form on twin-fault intersections in the material. The carbides precipitation occurs by mechanisms described by classical nucleation theory. • There was no relation between strain-induced martensite and TG carbide precipitation observed in this work. Acknowledgements--The authors would like to acknowl-
edge support received from the National Science Foundation (RIMI), through grant HRD-9105065, and from
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ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
EPA Cooperative agreement CR-818296-01-0 through the Southwest Center for Environmental Research and Policy. The work was also funded in part by GSA grant PF90-018, administered by the University of Texas at E1 Paso's Materials Research Institute. We are also extremely grateful to Dr Karl P. Staudhammer of Los Alamos National Laboratory for making the Ferritescope measurements reported in Fig. 7. REFERENCES
1. E. C. Bain, R. H. Aborn and J. J. B. Rutherford, Trans. Am. Soc, Steel Treating, 21, 481 (1933). 2. L. E. Murr, A. H. Advani, D. G. Atteridge and S. Shankar, Mater. Charact. 24, 135 (1990). 3. A. H. Advani, L. E. Murr, D. G. Atteridge and R. Chelakara, Metall. Trans. 22A, 2917 (1991). 4. A. H. Advani, D. G. Atteridge, L. E. Murr and R. Chelakara, Scripta metall, mater. 25, 343 (1991). 5. A. H. Advani, D. G. Atteridge, L. E. Murr, S. M. Bruemmer and R. Chelakara, Scripta metall, mater. 25, 461 (1991). 6. C. S. Tedmon Jr, D. A. Vermilyea and D. E. Bruecker, Corrosion 27, 104 (1971). 7. C. L. Briant and A. M. Ritter, Scripta metall. 13, 177 (1979). 8. C. L. Briant and A. M. Ritter, Metall. Trans. I1A, 2009 (1980). 9. C. L. Briant and A. M. Ritter, Metall. Trans. llA, 910 (1981). 10. V. Cihal, Corros. Sci. 20, 737 (1980). 11. V. Cihal, Intergranular Corrosion of Steels and Alloys, [pp. 146-157.] Elsevier Science Publishers, Amsterdam. (1984). 12. A. H. Advani, D. G. Atteridge and S. M. Bruemmer, Proc. Conf. on Precipitation Phenomena--Deformation and Aging Symposium, World Mater. Congress of Am. Soc. Metals, pp. 3-17 (1988). 13. A. H. Advani, D. G. Atteridge, L. E. Murr and S. M. Bruemmer, Proc. 2nd Int. Conf. on Trends in Welding Research (edited by S. A. David and J. M. Vitek), pp. 315-319. Am. Soc. Metals, Metals Park, Ohio (1990). 14. S. Pednekar and S. Smialowska, Corrosion 36, 565 (1980). 15. A. Bose and P. K. De, Corrosion 43, 624 (1987). 16. G. Rondelli, B. Mazza, T. Pastore and B. Vicentini, Electrochemieal Methods in Corrosion Research, (edited by M. Dupont), Mater. Sei. Forum 8, 593 (1986). 17. J. W. Simmons, D. G. Atteridge, L. E. Murr and S. M. Bruemmer, Proe. 2nd Int. Conf. on Trends in Welding
18. 19. 20. 21. 22. 23. 24. 25. 26. 27. 28. 29. 30. 31. 32. 33. 34. 35. 36. 37. 38. 39. 40. 41.
Research (edited by S. A. David and J. M. Vitek), pp. 321-325. Am. Soc. Metals., Metals Park, Ohio (1990). A. H. Advani and L. E. Murr, Scripta metall, mater. 25, 349 (1991). A. H. Advani, L. E. Murr, D. G. Atteridge, R. Chelakara and S. M. Bruemmer, Corrosion 47, 939 (1991). B. Weiss and R. Stickler, Metall. Trans. 3A, 851 (1972). N. Terao and B. Sasmal, Metallography 13, 117 (1980). C. Strawstrom and M. Hillert, J. Iron Steel Inst. 207, 77 (1969). L. P. Stoter, Z Mater. Sci. 16, 1039 (1981). J. K. L. Lai, Mater. Sei. Eng. 58, 195 (1983). M. H. Lewis and B. Hattersley, Acta metall. 61, 1159 (1983). L. K. Singhal, Metall. Trans. 12A 1267 (1971). L. K. Singhal, and J. W. Martin, Trans. metall. Soc. A.LM.E. 242, 814 (1968). L. K. Singhal and J. W. Martin, Acta metall. 15, 1603 (1967). L. K. Singhal and J. W. Martin, Acta metall. 16, 1159 (1968). A. Bouerf, R. Coppola, P. Zambonadi, J. P. Morlevat, F. Rustichelli and D. Wenger, J. Mater. Sci. 16, 1975 (1981). D. G. Morris and D. R. Harries, Metal Sci. 12, 542 (1978). G. B. Olson and M. Cohen, J. less-common Metals 28, 107 (1972). G. B. Olson and M. Cohen, Metall. Trans. 6A, 791 (1975). L. E. Murr, K. P. Staudhammer and S. S. Hecker, Metall. Trans. 13A, 627 (1982). K. P. Staudhammer, L. E. Murr and S. S. Hecker, Acta metall. 31, 267 (1983). K. P. Staudhammer, L. E. Murr and S. S. Hecker, Proc. Int. Conf. on Solid Phase Transformations, Carnegie Mellon Univ. (edited by H. Aaronson), p. 1287 (1983). S. S. Hecker, M. G. Stout, K. P. Staudhammer and J. L. Smith, Metall. Trans. 13A, 619 (1982). L. E. Murr and K. P. Staudhammer, Scripta metall. 16, 713 (1982). R. Stickler and A. Vinckier, Trans. Am. Soc. Metals 54, 362 (1961). R. Yong-Hua, G. Yong-Xiang, H. G. Xiang, Metallography 22, 47 (1989). D. A. Porter and K. E. Easterling, Phase Transformations in Metals and Alloys, pp. 142-147. Chapman and Hall, University and Professional Division, New York (1981).
0956-7151/93 $6.00.4-0.00 Copyright © 1993Pergamon Press Ltd
Printed in Great Britain. All rights reserved
DEFORMATION-INDUCED MICROSTRUCTURE A N D MARTENSITE EFFECTS ON TRANSGRANULAR CARBIDE PRECIPITATION IN TYPE 304 STAINLESS STEELS A. H. ADVANI, L. E. MURR, D. J. MATLOCK, R. J. ROMERO, W. W. FISHER, P. M. TARIN, J. G. MALDONADO, C. M. CEDILLO, R. L. MILLER and E. A. TRILLO Department of Metallurgical and Materials Engineering and Materials Research Institute, University of Texas at El Paso, E1 Paso, TX 79968, U.S.A. (Received 18 September 1992; in revised form 5 March 1993)
Abstract--Plastic deformation of 304 stainless steel (SS) induces transgranular (TG) carbide precipitation, which is critically dependent on deformation-induced microstructural changes occurring during thermal treatment of the SS. Uniaxial deformation of the 304 SS to 40% strain produces a high density of intersecting micro-shear bands composed of heterogeneous bundles of twin-faults and about 12-17% strain-induced ct'-martensite at the intersections of the twin-faults. Thermal treatment at 670°C for 0.1-10 h, however, results in a rapid annihilation/transformation of the strain-induced martensite and the concurrent formation of zones containing mixed thermal martensite laths and fine-grained austenite, though the thermal martensite also decreases with increasing heat treatment time. Simultaneous with these thermomechanically-induced microstructural changes, TG chromium-rich carbides form at intersections of twin-faults and on fine-austenite or thermal martensite boundaries in the SS; however, no correlation between strain-induced ~'-martensite and carbides was observed in this work. The mechanisms of deformation-induced microstructure and (strain-induced and thermal) martensite effects on TG carbide precipitation in 304 SS are discussed.
INTRODUCTION Sensitization refers to the development of grain boundary chromium-depletion caused by the precipitation of chromium-rich carbides on grain interfaces of stainless steel (SS) [1]. Sensitization occurs during thermal or thermomechanical (TM) treatment of SS, especially when the material is isothermally aged or continuously cooled through the temperature range of 500-850°C. Chromium-rich carbides precipitate on the grain boundary during these TM treatments, and due to a relatively slow diffusivity of chromium vs carbon, a substantial amount of chromium is drawn from the region around the growing carbide. This leads to grain boundary chromium-depletion, or sensitization, in the SS and renders the material susceptible to intergranular (IG) corrosion and IG stress-corrosion (SCC) cracking in certain aqueous environments. Example environments where SS is susceptible to IG corrosion and IGSCC include high temperature oxygenated water used in boiling water reactors, or corrosive environments present in containers used to store and transport toxic and radioactive waste materials where the temperatures can rise above ambient [2, 3]. Plastic deformation accelerates the development of grain boundary carbide precipitation and sensitization in 304 and 316 SS [3-17]. In recent work on 316 SS, we have shown that the acceleration of sensitization is a function of the amount of (prior) strain in the material and temperature of isothermal sensitiz-
ation treatment [3-5]. A systematic increase in strain from 0 to 20% was observed to yield a continuous increase in sensitization below 700°C, while at higher temperatures, a threshold strain of 5-10% was required to produce accelerated chromium depletion. Strain-induced acceleration in sensitization was also indicated to be caused by an acceleration in the kinetics of chromium diffusion, and by a lowering of the free energy barrier to carbide nucleation in the SS [3, 5]. Straining above 20% produces higher amounts of chromium depletion through the simultaneous precipitation of IG and transgranular (TG) (or grain matrix) carbides in 304 and 316 SS [12-19]. Transgranular chromium-depletion in 316 SS has been observed to be strain and heat treatment dependent, yielding C-curve behavior on a timetemperature-strain-TG corrosion map [18, 19]. The strain effect dominates, and increased amounts of TG depletion, and decreased time to develop TG corrosion occur with increasing strain level. Higher temperatures also reduce the time to develop grain matrix attack, while isothermal holding for longer times produces higher amounts of T G corrosion in SS. Transgranular depletion in 316 SS has been shown to occur due to precipitation of carbides within grain-matrix regions [18, 19]. Specifically, defect sites created during straining are favored locations for T G carbide precipitation in SS [19-31]. Transmission electron microscopy (TEM) has revealed that regions of high dislocation density,
2589
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ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
deformation twins, stacking-faults, and twin-fault intersections are preferred sites for matrix precipitation in deformed SS. Briant and Ritter [7-9] have compared the effects of strain on sensitization in 304 and 316 SS deformed to 5 ksi below the failure stress of the material, and indicated fundamental differences in deformation-induced sensitization for these materials. Briant and Ritter showed that 304 SS had accelerated T G and IG corrosion over 316 SS, and the amount of T G attack was especially more pronounced in 304 vs 316 SS. This was attributed to rapid precipitation of carbides, and development of chromium-depletion, on straininduced martensite and martensite lath boundaries in 304 SS, vs on slip planes in 316 SS. However, the evidence provided to relate the martensite to carbide precipitation only showed carbides on martensite laths, and the relationship between strain-induced martensite and carbide precipitation was not documented. To our knowledge, the evolution of deformation microstructure, martensite transformation, and T G carbide precipitation with TM treatment has also not been defined in any work to-date. Since the mechanisms and crystallography of strain-induced martensite formation in 304 SS are known [32-37], we believed that an understanding of TM processing effects on deformation microstructure (especially martensite), and identification of key relationships between strain-induced martensite and T G carbide precipitation, would be critical in evaluating mechanisms of carbide precipitation and chromiumdepletion phenomena in SS. The objective of this work was to develop a fundamental understanding of deformation-induced microstructure and martensite effects on T G carbide precipitation in 304 SS. The work was designed to focus on using TEM to investigate the relationship between T G carbides, deformation-induced sites and strain-induced martensite in 304 SS. Bulk X-ray diffraction (XRD) and magnetic measurement techniques were used to support the observed changes in microstructure during TM treatment of 304 SS. EXPERIMENTAL TECHNIQUES Material, deformation processing and heat treatment The material used in this work was a high carbon, Type 304 austenitic stainless steel of the composition shown in Table 1. The material was received as a 6.25 mm thick plate in the mill processed condition, and solution annealed at I100°C for 1 h prior to deformation processing. The solution annealed material had large, equiaxed grains, with an average grain-size of 175 gm. The microstructure of the annealed material was also
%C 0.051
confirmed to be clean, with a relatively low-dislocation density and no carbide precipitation on grain boundaries, which made it an appropriate initial condition for the deformation and heat treatment experiments. The solution annealed plate was cut into strips of length 125mm, and width 12.5mm. The sample was marked at 12.5mm intervals, and the initial a r e a (Ai) was measured at the marked locations on the sample. The strips were pulled in uniaxial tension using an Instron universal testing machine. The final area at the marked locations on the sample was then measured (Af), and the true strain was calculated using the relation true strain = In (Ai/Af). In this work, samples were deformed to ~ 4 0 % true strain, which corresponds to an engineering strain of ~49%. The deformed samples were cut into sections of 12.5 x 12.5 mm, and then heat treated at 670°C for 0.1-10 h to examine the effects of strain and heat treatment on carbide precipitation in SS.
Examination o f deformation microstructure and carbide precipitation in S S The deformation microstructure and carbide precipitation in 40% true strain, 670°C/0-10 h samples of SS were examined using a Hitachi H-8000 scanning transmission electron microscope (STEM), operated at 200 kV. Conventional bright-field TEM, dark-field TEM, selected-area electron diffraction and microdiffraction techniques were used for identification of phases produced during the TM treatments. Limited energy dispersive X-ray spectroscopy (EDS) was also used to identify chromium-rich carbide phases in SS. The formation of martensite in deformed samples, and the effect of heat treatment on the martensite transformation were also examined using XRD and magnetic measurement techniques. The austenite-to-martensite transformation was qualitatively studied using X R D by examining and comparing the face-centered cubic (f.c.c.) austenite peaks vs the body-centered cubic (b.c.c.) martensite peaks in the X R D spectrum. Also, since the martensite product is magnetic, the austenite-to-martensite transformation is ferromagnetic and was readily detected by magnetic measurements. The magnetic permeability of 40% strain, 670°C, 0-10 h samples was measured with a commercially available Ferritescope, which is calibrated to measure the content of ferrite in weldments. Additional details of the magnetic measurement technique to determine% martensite in SS are described elsewhere [37].
Table 1. Chemicalcompositionof type 304 stainlesssteel heat (in wt%) %Mn %P %S %Si %Ni %Cr %Mo %Co %Cu 1.47 0.025 0.014 0.38 8.22 18.29 0.19 0.11 0.19
%N 0.085
ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL RESULTS Plastic deformation effects on microstructural development
Plastic deformation to 40% strain produced several microstructural features in 304 SS, the most significant of which are illustrated in Fig. 1. The figure shows a bright-field TEM micrograph, a dark-field TEM micrograph and a microdiffraction pattern which illustrate the formation of strain-induced ~'martensite at twin-fault intersections in the as-deformed SS. The bright-field micrograph indicates the presence of intersecting microshear bands in the deformed microstructure of the 304 SS. These microshear bands are usually composed of heterogeneous bundles of stacking faults and deformation twins, and in some cases may also include E-martensite. Straininduced martensite is seen to form within the twinfault intersections; locations of the martensite are clearly indicated by the arrows in the bright field micrograph, and also by the "lit-up" areas in the dark-field micrograph. The presence of strain-induced martensite was confirmed by microdiffraction analysis of the twin-fault intersections which contained martensite.
i1o
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BCC c~' (I11) Fig. 1. Illustration of strain-induced martensite formation in 304 SS. The TEM micrographs are bright-field and darkfield views of the strain-induced ~'-martensite at twin-fault intersections in the SS. The arrows in the bright-field view show some locations of the ~'-martensite in the micrograph. The diffraction pattern results by microdiffraction from a martensitic intersection. AMM 41 t~-F
2591
Only a fraction of the twin-fault intersections in the SS was observed to transform to strain-induced martensite. Bulk ferromagnetic analysis, in fact, quantified the fraction of ~'-martensite in the 40% strain, as-deformed 304 SS to be 12-17% (depending on the orientation of measurement). X-ray diffraction analysis also qualitatively confirmed that a fraction of the 304 SS material was transformed to strain-induced martensite, as can be seen in a comparison of relative peak heights for ~'-martensite and y-austenite peaks shown in Fig. 2. Thermal treatment effects on defi~rmation microstructure
Heat treatment of the 40% true strain 304 SS produced significant and rapid changes in the deformation microstructure of the material. For the shortest time of heat treatment examined, i.e. 0.1h at 670°C, no strain-induced martensite was observed at twin-fault intersections, unlike that seen in the asstrained material. Instead, a "new" cell-like phase was seen to form at the twin-fault intersections, as is shown in the bright-field micrographs and microdiffraction patterns of Fig. 3(a) and (b). The high magnification view of Fig. 3(a) illustrates the elongated cellular morphology of the phase. Analysis of the microdiffraction of Fig. 3(b) indicated that the phase was b.c.c.-martensite. However, the cell-like appearance of the phase, and the extent of its growth even outside twin-fault intersections suggested that the phase was not strain-induced martensite, but was more typical of lath martensite; hence we considered the phase to be "newly" formed during the heat treatment. It was also apparent that the lath martensite formed during the heat treatment process, while the strain-induced martensite was seen only in (room temperature) strained SS; hence, we considered the lath martensite to be a "thermally" induced, rather than a strain-induced, product. And, to assist in our differentiation between the two martensites, we have chosen to use the terms "lath martensite" and "thermal martensite" to refer to the thermally-induced phase, while the "strain-induced-martensite" and "~'-martensite" will continue to imply the strain-induced phase. Other locations which appeared to be formed by a continued production of the thermal martensite were also observed in this 670-C-0.1 h sample (Fig. 4). The cell-like regions again originated within the twin-fault intersection regions, as is illustrated in the bright-field and dark-field micrographs of Fig. 4. Selected-area electron diffraction indicated a complex ring pattern, which is possibly due to the fine-intermix of thermal martensite dispersed within the austenite at the location. Longer heat treatments to 0.4 h at 670'~C resulted in an increased observation of the clustered bands of cellular phases, and in several locations the clusters often encompassed entire sections of grains (Fig. 5). The clustered cells appeared to be oriented in specific
2592
ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 2. X-ray diffraction spectrum showing martensite and austenite peaks in the spectrum. The ~t' represents the strain-induced martensitic peaks, while the ? represents the austenitic peaks. directions, and often appeared as bands in the material. Although some clusters were also observed in the 40% strain, 670°C-0.1 h sample, the extent of clustering was more prominent in the 0.4 h sample. Selected-area electron diffraction again indicated the presence of a complex ring-pattern in the region. Further microdiffraction analysis was, thus, essential within the clusters in this sample to confirm the phases within the cell-locations. Microdiffraction analysis of the clustered region indicated that the cellular region was not entirely thermal martensite, but instead contained a mixture of thermal martensite and fine-grained austenite (Fig. 6). We will refer to this fine-grained austenite that forms within local, T G clustered regions as "fine-austenite" to distinguish it from the largegrained austenite in the initial microstructure. An example dimension of the fine-austenite grain-size was measured to be 0.2 x 0.3 #m, in comparison with the original solution-annealed grain size of 175 #m. Both the fine-austenite and lath martensite are seen to have relatively low-dislocation densities in Fig. 6. The fine-austenite may also be contributing to the complexity of ring selected-area electron diffraction patterns observed in Figs 4 and 5. Additional heat treatment to 1 and 10 h at 670°C continued to show the presence of clustered cellular regions in the microstructure. No significant changes in the fine-austenite size or lath martensite size were seen from the 0.1-I0 h of treatment. Microdiffraction
of several locations again indicated that both fineaustenite and martensite were present within the clusters, although it was considered possible that there may be changes in the amount of martensite vs austenite ratios in these samples. To verify this possibility, we used bulk magnetic measurements and X R D analysis of all heat treated samples. Ferromagnetic analysis indicated that the amount of martensite changed dramatically with longer heat treatments at 670°C. The martensite was relatively constant (about 12-17%) from 0.0 to 0.4 h of heat treatment, and then dropped sharply to 2-3% after 1 and 10 h at 670°C (Fig. 7). The detection of the martensite by magnetic techniques was also seen to have an orientation effect, which is not understood. XRD analysis supported the bulk magnetic observations and indicated the same qualitative trends in martensite level and an orientation effect of the martensite, though it was difficult to quantify the martensite during X R D (Fig. 8). Thus, both these bulk techniques indicate that the clustered "cellular" regions seen in the 1 and 10 h samples are almost always austenite, and our observations of some lath martensite in the TEM samples is probably a statistical effect. Thermal treatment effects on transgranular carbide precipitation
Heat treatment also resulted in T G carbide precipitation, which occurred simultaneously with the
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 3. Formation of thermal martensite laths at twin-fault intersections in the 40% -670°C4). 1 h sample.
Fig. 4. Formation of clustered regions at twin-fault intersections in the 40%-670°C~0.4h sample. Micrographs represent bright-field and dark-field views of the clusters. The selected-area electron diffraction pattern shows a complex ring diffraction pattern is obtained from the location.
Fig. 5. Illustration of the enhancement of the cluster~l region in the 40% strained sample with 0.4 h of treatment at 670°C. The clusters are seen to have a specific orientation, and the diffraction pattern is complex showing a ring pattern (and texturing effects) due to the fine intermixing of thermal lath martensite and untransformed austenite.
2594
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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deformation-induced microstructural changes in the SS. It should be noted that while carbides were also observed to form on grain boundaries in the 40% deformed SS, we have centered on the understanding of T G carbide precipitation in the SS, which is the focus of this work. Two primary locations for T G carbide precipitation were observed in this work: (1) on the boundaries within the mixed fine-austenite and martensite lath regions; and (2) on intersections of twin faults in the SS. An example view of an isolated carbide at a region of fine-austenite and martensite in the 0.4h sample is shown in Fig. 9. It illustrates the complexity of determining the presence of T G carbides in these materials, and indicates the key role of microdiffraction in the analysis• The brightfield TEM micrograph and corresponding selectedarea electron diffraction pattern did not provide a clear indication of the phases present at the location• However, microdiffraction at locations 1, 2 and 3 clearly reveals the presence of austenite, lath martensite, and carbides at these locations, respectively. It can be seen that the fine-austenite and martensite regions are completely intermixed. The carbide seems to form at a martensite lath boundary/fineaustenite boundary within the region• Additional evidence for the formation of carbides on the boundaries of mixed, fine-austenite grains and thermal martensite laths in these materials is provided in Fig. 10.
In the samples heat treated beyond 1 h and up to 10 h at 670°C (where only some thermal martensite was present), carbides were continued to be observed within the clustered regions. This is illustrated in the TEM micrographs showing lattice images of carbides formed in a predominantly fine-austenitic region (Fig. 11). Detailed views of the lattice images show undistorted carbide lattices, dislocations at the carbide-matrix interface, and changes in lattice orientation across the fine-austenitic interface. The other important location for T G carbides precipitation was on twin-fault intersections in the SS. Transgranular carbides on the intersections were difficult to observe, especially in the short heat-treatment (0. I, 0.4 h) samples• This is probably due to a small size of carbides, low density of carbides, or due to a well-spread distribution of carbides within the grain• Larger carbides were, however, also seen in the longer heat treatment (670°C-10 h) samples• Examples of small carbides on twin-fault intersections in the 0.4 h sample are shown in Fig. 12, while larger carbides seen at twin-fault intersections in the 670°C-10h sample are illustrated in Fig. 13. No strain-induced martensite was seen in the vicinity of the carbides in these or other twin-fault locations of the SS. DISCUSSION
Plastic deformation produces strain-induced ~t'martensite formation at the intersections of twin-faults in as-strained 304 SS. Subsequent heat treatment results in an annihilation/transformation of the strain-induced martensite, and in the simultaneous nucleation of thermal lath martensite, though the lath martensite also eventually decreases with increasing heat treatment time. The strain-induced martensite forms via the specific 20
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Fig. 7. Ferromagnetic analysis indicating variations in the amount of martensite with heat treatment. It should be noted that the amount of martensite measured was dependent on the orientation of the sample (face or edge) on which the martensite was recorded. The low martensite values of 2-3% should also be noted to be inaccurate as they a r e typically below the detection limit of the instrument. (Data courtesy Dr K. P. Staudhammer, Los Alamos National Laboratory.)
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3C6
Fig. 9. An example view of an isolated TG carbide at a region of fine-austenite and thermal martensite in the 40%--670°C-0.4 h sample. The bright-field image is on the top-left of the figure and a selected-area electron diffraction taken from the entire region is to the right of the bright-field. Microdiffraction patterns for fine-austenite, thermal martensite and TG carbide (bottom left to right) were taken from locations I, 2 and 3 marked on the bright-field micrograph respectively.
crystallographic conditions defined by Olson and Cohen for 304 SS [32, 33]. The volume of strain-induced martensite formed is also strain-dependent, and the 12-17% transformation observed in this work, reflects the 40% strain in the material [37]. The requirement for a thermal impetus to form the lath martensite has not been previously documented in deformed SS, although some changes in martensite morphology were reported by Murr and Staudhammer in shock-loaded 304 SS [38]. And although the lath martensite and strain-induced martensite phases have the same b.c.c, crystal structure and lattice parameter in SS, we believe that the mechanisms and crystallography of transformation for the two martensites would be significantly different. While no direct evidence of a relationship between strain-induced ~t'-martensite and the lath martensite was noted, the results allude to a possible transformation from strain-induced martensite to lath martensite, for the following reasons. First, both phases formed on twin-fault intersections in the SS, and the observation of the lath martensite coincided with the disappearance of the strain-induced martensite, all of which occurred within a short duration of 0.1 hr at 670°C. And second, the bulk observations also sup-
port the possibility of a complete transformation from strain-induced martensite to thermal lath martensite, This is because the total volume of martensite was essentially unchanged between the as-deformed and 670°C-0.1 and 0.4 h samples, and only thermal martensite was seen in the heat treated samples while only strain-induced martensite was seen in the as-deformed samples. However, the evidence is not conclusive, and it is also possible that the strain-induced martensite is annihilated at twin-fault intersections during the heat treatment, while the lath martensite forms at other unrelated twin-fault intersections within the same period. The clustered regions are produced by a fine intermixing of thermal martensite laths with fine-austenite, and may also include untransformed austenite grains. The clusters form within 0.1 h of heat treatment, though the thermal lath martensite eventually decreases after 1 h at 670°C. The formation of the fine-austenite could be caused by a local-recrystallization process, simultaneous transformation of straininduced or lath martensite to fine-austenite, or local encompassment of untransformed austenite by the thermal martensite laths. And, the lack of growth within fine-austenite grains observed with heat
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
iC iii!!! iiiii(!i
Fig. 10. Formation of TG carbides on the boundaries of mixed, fine-austenite grains and thermal martensite laths in 304 SS.
treatment may be because the 670°C temperature does not provide sufficient thermal impetus for the growth of these regions, or because the grains are pinned by carbides which form on their boundaries. It should be noted that the T G clustered regions discussed above are similar to the lath martensite regions previously described by Briant and Ritter [7-9]. However, in this work, we have provided a clearer definition of the formation of the region, and conclusively demonstrated that it consists of a fine intermix of thermal lath martensite and fine-austenite, rather than only lath martensite (or strain-induced martensite) as was described previously. Transgranular carbide precipitation is critically dependent on the formation of the clustered regions in the SS, which matches similar observations made by Briant and Ritter [7-9]. The formation of T G carbides on boundaries within the mixed fine-austenite and martensite lath regions is a reflector of the high energetics associated with ferrite-austenite interfaces; in fact, Stickler and Vinckier [39] have shown that carbides prefer to precipitate first on ferrite-austenite boundaries, in comparison with grain boundaries, twin-boundaries, and other deformation sites in SS. However, the loss of thermal martensite after 1 h of heat treatment does not seem to influence the T G carbide precipitation, since the precipitation occurs even at the longer heat treatments; the precipitation on the now completely fine-
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austenite regions would be similar to carbide formation on grain boundaries in SS. Thus, while the thermal martensite is a contributor to the TG carbide precipitation process, it is not a necessary requirement for the precipitation to occur within the clustered zones in the SS. The observed lattice images of T G carbides within the predominantly fine-austenitic cluster zone observed in the 670°C-1 h sample indicate key mechanisms of carbide precipitation on grain boundaries in SS. The lattice changes associated with the carbide precipitation in our work on 304 SS have also been reported by Young-Hua et al. [40] during precipitation of M23C6 in a superalloy. The undistorted lattice image that was observed is indicative of a possible coherent match-up between the matrix and precipitate, in which the fringes may pass the carbide-austenite interface undisturbed into the austenite matrix; although no lattice fringes were observed in the matrix of our samples, Yong-Hua et al. have observed this result, and attributed it to the lattice spacing of M23C6 being approximately three times the spacing of the matrix. The dislocations seen on the carbide-austenite interface are due to the carbide trying to minimize the free energy of transformation by maintaining a semi-coherent interface with the matrix, in which the disregistery is taken up by misfit dislocations at the interface [41]; the occasional presence of misfit dislocations was also noted in YongHua's work. Furthermore, the change in lattice orientation across the interface has been suggested [40] to be caused by the carbide nucleating separately on either side of the grain boundary, so as to maintain a high degree of coherency with the grain in which it grows. Thus, there are a range of classical precipitation mechanisms by which the carbide grows in 304 SS, and we believe that the preferred mechanism is specific to each precipitate, so as to minimize the free energy of the system. Transgranular carbides also formed on twin-fault intersections in the 304 SS. However, no evidence of any relation between strain-induced martensite and T G carbides was observed in this work, except that they both preferred to form on twin-fault intersections in the SS. This observation is in contradiction with previous results reported by Briant and Ritter [7-9], and the differences could be due to heat-to-heat variations in SS, TM processing differences, or limited area observed during TEM. While there may or may not be a correlation between strain-induced martensite and T G carbides, it is clear that the martensite phase is definitely not a precursor to the formation of carbides on twin-fault locations in SS. This is supported by our observation of carbide precipitates at twin-fault locations where there was no strain-induced martensite. The absence of a requirement for strain-induced ~'-martensite to be a precursor to the formation of T G carbides is further supported by the fact that carbides also form on twin-fault intersections in 316 SS [18, 19], where the
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CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
Fig. 1I. TEM micrograph showing lattice images of TG carbides formed in a fine-austenitic region. Detailed views of the lattice images, in fact, show the undistorted carbide lattice, dislocations at the carbide-matrix interface, and changes in lattice orientation across the fine-austenitic interface.
~53
353 FCC-M23C6 ( 2-37-)
5h
220
i.~sT 511 FCC-M23C6 (116)
Fig. 12. Example views of TG carbides on twin faults and their intersections in the 670°C4).4 h sample (a, b). No strain-induced martensite was seen to be present in the vicinity of the carbide.
ADVANI et al.:
CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
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Fig. 13. Illustration of carbides seen at TG twin-fault intersections in the 670°C-10 h sample (a) and (b).
strain-induced martensite transformation does not Occur.
Thus, the key mechanisms that define the role of deformation-induced microstructure and martensite formation on TG carbide precipitation have been identified and fundamentally evaluated in this work. It is clear that it is not only plastic deformation, but the synergistic effect of the deformation and thermal treatment, which produce the required microstructural sites that are key to the precipitation of the T G carbides in the SS. The strain-induced martensite phase which lbrms in the as-deformed material is annihilated/transformed during the thermal treatment and, therefore, has no correlation with the carbide precipitation in SS. Instead clustered zones that consist of thermal lath martensite and fineaustenite, and twin-fault intersections (with no martensite), are more potent sites for the T G carbide formation, which occurs by classical nucleation mechanisms to minimize the free-energy of the system. And, while the lath martensite contributes to the TG precipitation process, neither the strain-induced martensite nor lath martensite are a precursor or required for carbide precipitation in the SS. All these evidences point to the fact that the energetics, rather than the crystal structure, of the site is a key to the precipitation of TG carbides in the SS. Ultimately, the energetically driven, deformation site-specific behavior of TG carbides will control the sensitization,
corrosion and SCC behavior of heavily deformed SS in service applications. SUMMARY AND CONCLUSIONS QPlastic deformation of 304 SS to 40% strain produces microshear bands composed of twin-faults and about 12-17% strain-induced martensite. QThermal treatment of the deformed SS results in the annihilation/transformation of the strain-induced martensite, and the simultaneous formation of mixed regions containing about 12-17% thermal martensite laths and fine-austenite. The lath martensite formation also occurs at twin-fault intersections in the SS. Longer heat treatment (beyond 670°C-1 h) results in a reduction in the amount of thermal martensite in the SS to 2-3%. • Transfer anular carbides form on the mixed, thermal martensite lath and fine-austenite boundaries in the SS. They also form on twin-fault intersections in the material. The carbides precipitation occurs by mechanisms described by classical nucleation theory. • There was no relation between strain-induced martensite and TG carbide precipitation observed in this work. Acknowledgements--The authors would like to acknowl-
edge support received from the National Science Foundation (RIMI), through grant HRD-9105065, and from
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ADVANI et al.: CARBIDE PRECIPITATION IN 304 STAINLESS STEEL
EPA Cooperative agreement CR-818296-01-0 through the Southwest Center for Environmental Research and Policy. The work was also funded in part by GSA grant PF90-018, administered by the University of Texas at E1 Paso's Materials Research Institute. We are also extremely grateful to Dr Karl P. Staudhammer of Los Alamos National Laboratory for making the Ferritescope measurements reported in Fig. 7. REFERENCES
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