An Investigation of the Massive Transformation from Ferrite to Austenite in Laser-Welded Mo-Bearing Stainless Steels

Transcription

1 An Investigation of the Massive Transformation from Ferrite to Austenite in Laser-Welded Mo-Bearing Stainless Steels M.J. PERRICONE, J.N. DuPONT, T.D. ANDERSON, C.V. ROBINO, and J.R. MICHAEL A series of 31 Mo-bearing stainless steel compositions with Mo contents ranging from 0 to 10 wt pct and exhibiting primary d-ferrite solidification were analyzed over a range of laser welding conditions to evaluate the effect of composition and cooling rate on the solid-state transformation to c-austenite. Alloys exhibiting this microstructural development sequence are of particular interest to the welding community because of their reduced susceptibility to solidification cracking and the potential reduction of microsegregation (which can affect corrosion resistance), all while harnessing the high toughness of c-austenite. Alloys were created using the arc button melting process, and laser welds were prepared on each alloy at constant power and travel speeds ranging from 4.2 to 42 mm/s. The cooling rates of these processes were estimated to range from 10 K ( C)/s for arc buttons to 10 5 K( C)/s for the fastest laser welds. No shift in solidification mode from primary d-ferrite to primary c-austenite was observed in the range of compositions or welding conditions studied. Metastable microstructural features were observed in many laser weld fusion zones, as well as a massive transformation from d-ferrite to c-austenite. Evidence of epitaxial massive growth without nucleation was also found when intercellular c-austenite was already present from a solidification reaction. The resulting singlephase c-austenite in both cases exhibited a homogenous distribution of Mo, Cr, Ni, and Fe at nominal levels. DOI: /s x Ó The Minerals, Metals & Materials Society and ASM International 2010 I. INTRODUCTION THE massive phase transformation is a compositioninvariant, interface-controlled diffusional phase transformation. [1] Sometimes referred to as partitionless, [2] massive transformations require only short-range diffusion across the massive/parent interface to change one phase to another. [1 10] The cooling rate must be sufficiently high to quench the material into a two-phase field below T 0 while suppressing competing diffusional transformation mechanisms that can occur above T 0, thereby accumulating sufficient driving force to facilitate this mechanism. Massive transformations have been observed in multiple materials systems [1 10] and growth rates exceeding 10 mm/s have been reported, [1] though much lower speeds are also possible. [4,5] The absence of long-range diffusion causes the massive product phase to inherit the composition and chemical distribution of the M.J. PERRICONE, formerly Senior Member Technical Staff, Sandia National Laboratories, Albuquerque, NM 87185, is Senior Scientist, RJ Lee Group, Inc., Monroeville, PA Contact mperricone@rjlg.com J.N. DuPONT, Professor, is with the Department of Materials Science and Engineering, Lehigh University, Bethlehem, PA T.D. ANDERSON, formerly Graduate Research Assistant, Department of Materials Science and Engineering, Lehigh University, is Senior Research Engineer, Exxon Mobil Corporation, Irving, TX C.V. ROBINO, Distinguished Member Technical Staff, is with the Joining and Coatings Division, Sandia National Laboratories. J.R. MICHAEL, Distinguished Member Technical Staff, is with the Materials Characterization Department, Sandia National Laboratories. Manuscript submitted June 19, Article published online November 19, 2010 parent, a characteristic useful for identifying phases formed by this phenomenon. The compositional invariance that distinguishes massive transformations from other diffusional solid-state transformations may be particularly advantageous in austenitic stainless steels, which rely on the distribution of critical alloying elements for their corrosion resistance. Weldable commercial austenitic stainless steels often have compositions that solidify as primary ferrite [11] to avoid the solidification cracking susceptibility typical of primary austenite solidification, [12] followed by a solid-state transformation to austenite. However, as current trends continue to increase alloy content (e.g., superaustenitics and superduplexes) to maximize corrosion resistance in aggressive environments, the challenges associated with welding such alloys becomes a major source of concern. It is well known that alloying additions can change the primary solidification mode for a given alloy, and the non-uniform redistribution of critical alloying elements (especially Mo) during primary austenite solidification [13 15] results in depletion in localized regions of the fusion zone microstructure, leaving them susceptible to preferential corrosive attack. [16 19] Furthermore, the concomitant localized buildup of solute in other areas of the microstructure can promote the formation of brittle intermetallics such as sigma. [20] While primary d-ferrite solidification can ameliorate these effects to some extent (higher solubility for Mo, tramp elements, etc.), many applications require the high toughness and lack of magnetic signature afforded by austenite. Recent work by Anderson et al. [21] 700 VOLUME 42A, MARCH 2011

2 demonstrated the existence of a broad range of high alloy stainless steel compositions that can simultaneously capture the immunity to solidification cracking via primary ferrite solidification and the room-temperature properties of austenite via a solid-state transformation from d-ferrite. The present investigation focuses on the solid-state transformation from d-ferrite to c-austenite in laser welds prepared on alloys in the Fe-Ni-Cr-Mo system. The results of this investigation will be viewed in the context of several studies [11,22 24] that have reported a massive transformation from d-ferrite to c-austenite in high energy density (HED) welds prepared on Fe-Ni-Cr ternary alloys. A similar investigation of this phenomenon in the higher order alloy systems on which most engineering stainless steels are based is absent from the literature, as are potential avenues through which this transformation mechanism can be controlled. Because the microstructural development of this class of materials has been shown to rely heavily on the redistribution of sluggishly diffusing elements such as Mo, clarification of the role of composition and cooling rate in massive transformation phenomena that can occur in HED welds in Mo-bearing stainless steel alloys is warranted. II. EXPERIMENTAL PROCEDURE Target alloy compositions (Table I) were chosen based on a series of thermodynamic simulations presented elsewhere [21,25] and were prepared from highpurity materials using a button arc melting process in heats of approximately 50 g. Average compositions were measured using inductively coupled plasma (ICP) wet chemical analysis, the results of which are also presented in Table I. A subset of these alloys was also evaluated by using combustion analysis to determine C, O, and N levels (Table II). Autogenous laser welds were prepared on the bottom of each button using a 500 W continuous wave Nd-YAG laser in a chamber containing a pure Ar atmosphere under pressure slightly above atmospheric to prevent ingress of air. Power at the button surface was measured at 370 W using a puck -type calorimeter. Travel speeds of 4.2, 21, and 42 mm/s were used to provide a range of cooling rates and solidification velocities. Welds were made ~5 mm apart from each other to avoid overlap of the temperature fields associated with each pass. Samples were sectioned in the transverse direction and were prepared for microstructural observation using standard metallographic techniques. This included grinding with successively Table I. Results of ICP Wet Chemical Analysis; the Balance of All Alloys is Fe; All Values are Expressed in Weight Percent Sample Mo Cr Ni Fe* Al Cu Si Mn 0Mo20Cr10Ni < Mo22Cr13Ni < Mo22Cr8Ni <0.01 < Mo24Cr11Ni Mo18Cr12Ni Mo20Cr10Ni Mo22Cr13Ni Mo24Cr11Ni Mo17Cr13Ni Mo17Cr8Ni Mo19Cr11Ni Mo20Cr15Ni Mo22Cr13Ni <0.01 < Mo13Cr12Ni <0.01 < Mo15Cr10Ni < Mo16Cr14Ni <0.01 < Mo18Cr12Ni <0.01 < Mo20Cr15Ni <0.01 < Mo22Cr13Ni <0.01 < Mo13Cr12Ni Mo15Cr10Ni Mo15Cr15Ni Mo17Cr13Ni < <0.01 8Mo20Cr15Ni Mo22Cr13Ni Mo13Cr12Ni Mo14Cr16Ni Mo15Cr10Ni Mo16Cr14Ni < < Mo18Cr17Ni Mo20Cr15Ni *Determined by balance. VOLUME 42A, MARCH

3 finer grits of SiC, followed by polishing with 0.3-lm diamond paste, and a final polish on colloidal silica (SiO 2 ). The etchant used was electrolytic 60 pct HNO 3 / 40 pct H 2 O. Electron probe microanalysis (EPMA) was conducted on unetched specimens to evaluate the distribution of alloying elements within the fusion zone of each weld using a JEOL* 733 Superprobe equipped with four *JEOL is a trademark of Japan Electron Optics Ltd., Tokyo. independent wavelength dispersive spectrometers. An accelerating voltage of 15 kv and a beam current of 30 na were used. The K a lines were used for elements Fe, Ni, and Cr, while the L a line was used for Mo. [26] Raw data were converted to weight percentages through comparison with data collected from pure elemental standards and application of a phi(qz) correction scheme. Monte-Carlo simulations demonstrated that the interaction volume was 1 lm 3 for the conditions described here. Line scans were conducted across the fusion zone microstructure with a 2-lm step size between data points. To achieve higher spatial resolution in the refined scale of the fusion zone microstructures of the fastest laser welds in this study, transmission electron microscopy was employed to analyze primary c-austenite and primary d-ferrite compositions of laser welds prepared at 42 mm/s travel speed (~10 5 C/s cooling rate). Thin film specimens were prepared from specific portions of polished metallographic cross sections using the focused ion beam (FIB) milling technique on an FEI DB235 system and thinned to electron transparency in situ. The samples were then removed and placed on a carbon film for insertion into an FEI Technai F30 (FEI Worldwide Corporate Headquarters, Hillsboro, OR) equipped with an EDAX (EDAX 91, Mahwah, NJ) thin window spectrometer. Compositional mapping of 2.5-lm 2 regions for Fe, Ni, Cr, and Mo was accomplished using 200 kv accelerating voltage and a 9-nm step size. This spatial regime was chosen to ensure that an appreciable number of dendrites would be sampled (dendrite spacing Table II. Results of Testing for Carbon, Oxygen, and Nitrogen; All Values are Expressed in Weight Percent Sample C O N 4Mo20Cr15Ni < Mo13Cr12Ni Mo16Cr14Ni Mo20Cr15Ni Mo15Cr15Ni < Mo18Cr17Ni <0.001 ~1 lm) while still achieving good spatial resolution to capture any compositional gradients that may be present. Phase identification was conducted on regions of arc button and laser weld fusion zone microstructures in several alloys. Electron backscattered diffraction (EBSD) pattern analysis was used to positively identify the crystal structure of unknown phases without the need for thin film sample preparation. [27] Samples in the as-polished condition were examined in a Supra 55VP scanning electron microscope with the HKL Technology Channel 5 EBSD system (Oxford Instruments, Abingdon, Oxfordshire, UK). Imaging was accomplished using a KE Development forward scatter detector positioned low on the EBSD phosphor screen. [28] Microstructural imaging was also accomplished using band contrast, which is a measure of the relative contrast of the diffraction bands in a single EBSD pattern and can allow clearer delineation between grains within a phase than standard secondary or backscattered electron signals. The sample surface was tilted to 70 deg relative to the horizontal and analysis was completed at an accelerating voltage of 20 kv, a beam current of 2 to 3 na, and a working distance of 13 mm. The acquired patterns were indexed against published crystallographic data [29] to accurately distinguish between fcc (c-austenite) and bcc (d-ferrite and v-chi). The tetragonal (r-sigma) structure was compared to a Fe-Cr r-sigma phase prototype crystal structure. [30] The details of each crystal structure are summarized in Table III. Maps were created using a step size ranging from 0.05 to 0.2 lm depending on the minimum size of the features of interest in a given field of view. Mild levels of postprocessing were used to fill in unindexed pixels in the orientation maps shown here. III. RESULTS AND DISCUSSION The cooling rate (e, K( C)/s) for each solidification condition was estimated based on the measurement of primary dendrite arm spacing (k, lm). According to several authors, [23,31] the expression developed for 310 stainless steel by Katayama and Matsunawa, [32] k ¼ 80e 0:33 can be used to describe the behavior of Fe-Ni-Cr alloys with generally similar compositions, even to estimate cooling rates up to 10 6 K ( C)/s. [22,23] The results of these measurements and subsequent calculations are displayed in Table IV. The cooling rate was found to vary from 20 K ( C)/s in the arc buttons to 10 4 Kto Table III. Crystallographic Information Used to Identify Phases in the Fe-Ni-Cr-Mo Quaternary System Phase Name Crystal Structure Lattice Parameters Space Group Chemical Formula Reference Austenite fcc a = to nm Fm3m 53 Ferrite bcc a = to nm Im3m 53 Sigma tetragonal a = 0.879, c = P4 2 /mnm Fe-Cr-Mo 54 Chi bcc a = I43m Fe 36 Cr 12 Mo VOLUME 42A, MARCH 2011

4 10 5 K( C)/s in the laser welds. The primary solidification mode of each alloy in this study (Table V) was morphologically determined [12,23] to either be fully ferritic (F) or primary d-ferrite with eutectic c-austenite (FA). No secondary or tertiary dendrite arms were experimentally observed in the laser weld microstructures examined in this study. A. Primary Ferrite (F) Solidification The F mode of primary solidification is characterized by complete solidification of d-ferrite with no formation of secondary phases during solidification. In many of Melt Table IV. Dendrite Arm Spacing Measurements Used to Estimate Cooling Rates in this Study Travel Speed (mm/s) k (lm) Standard Deviation (lm) e, K( C)/s Button LW LW LW these alloys, a solid-state transformation to c-austenite is observed on cooling, exhibited by grain boundary allotriomorphs followed by the formation of Widmanstätten side plates extending into the parent d-ferrite grain. The details of this solidification mode and subsequent solid-state transformations under cooling conditions typical of conventional arc welding have recently been published by Anderson et al. [21] However, four alloys in the present study exhibited F primary solidification modes and transformed to single-phase c-austenite in the laser weld fusion zone: 2Mo20Cr10Ni, 2Mo22Cr13Ni, 6Mo15Cr10Ni, and 10Mo13Cr12Ni. For comparison, Figure 1(a) depicts a single-phase d-ferrite structure resulting from F mode solidification in a 4.2 mm/s weld prepared on alloy 8Mo17Cr13Ni. Note that the Widmansta tten c-austenite present in the base metal adjacent to the weld has been completely suppressed in the fusion zone by the high cooling rates associated with laser welding. Conversely, large blocky austenite grains, exhibiting flat and faceted c/d interfaces, are observed in a weld prepared using the same welding parameters on alloy 2Mo20Cr10Ni (Figure 1(b)). Regions of single-phase c-austenite were also observed to cross parent d-ferrite grain boundaries Table V. Primary Solidification Mode as a Function of Cr eq /Ni eq Ratio; Measured Compositions Were Used to Calculate Cr eq / Ni eq Ratios When Available; Italicized Primary Solidification Mode Indicates Conditions Where a Massive Transformation Occured from d-ferrite to c-austenite Sample Cr eq /Ni eq Mo Button 4.2 (mm/s) 21 (mm/s) 42 (mm/s) 10Mo14Cr16Ni FA FA FA FA 6Mo13Cr12Ni FA FA FA FA 6Mo16Cr14Ni FA FA FA FA 4Mo17Cr13Ni FA FA FA FA 2Mo18Cr12Ni FA FA FA FA 8Mo15Cr15Ni FA/F F*, ** FA FA 8Mo13Cr12Ni F FA FA FA 0Mo22Cr13Ni FA FA FA FA 4Mo20Cr15Ni FA FA FA FA 6Mo20Cr15Ni F F F F 10Mo18Cr17Ni FA/F F F F 10Mo16Cr14Ni F F F F 10Mo13Cr12Ni F F F F 8Mo20Cr15Ni F F F F 10Mo20Cr15Ni F F F F 2Mo22Cr13Ni FA/F F F F 8Mo17Cr13Ni F F F F 0Mo20Cr10Ni F FA FA FA 6Mo15Cr10Ni F F** F F 4Mo22Cr13Ni F F F F 6Mo18Cr12Ni F F F F 4Mo19Cr11Ni F F F F 0Mo24Cr11Ni F F F F 6Mo22Cr13Ni F F F F 2Mo20Cr10Ni F F F F 8Mo15Cr10Ni F F F F 2Mo24Cr11Ni F F F F 8Mo22Cr13Ni F F F F 10Mo15Cr10Ni F F F F 4Mo17Cr8Ni F F F F 0Mo22Cr8Ni F F F F *Local base metal microstructure was F. **Morphologically identified massive c-austenite. VOLUME 42A, MARCH

5 Fig. 1 Examples of metastable microstructures observed under laser welding conditions. (a) Single-phase d-ferrite in the fusion zone of 4.2 mm/ s weld prepared on F mode alloy 8Mo17Cr13Ni, (b) through (d) massive c-austenite morphology for laser weld prepared at 4.2 mm/s on F mode alloy 2Mo20Cr10Ni, (c) c m growth unimpeded by prior d grain boundaries, and (d) residual d located between c m that failed to grow together. unimpeded (Figure 1(c)), indicating the lack of an orientation relationship between c and the parent d-ferrite. Residual d-ferrite was observed in regions of the fusion zone microstructure, where c-austenite grains failed to fully impinge on each other, as shown in Figure 1(d). Note also the wavy appearance of the interface between the austenite and residual ferrite in some locations in Figure 1(d). The microstructures of F mode laser welds were subsequently examined using EBSD analysis to confirm phase identification and relative orientation. Figure 2 demonstrates the single-phase d-ferrite fusion zone that results from a laser weld prepared at 4.2 mm/s on alloy 10Mo18Cr17Ni, with only small amounts of c-austenite present. Of particular note in this analysis is the Kurdjumov Sachs (K-S) or Nishiyama Wasserman (N-W) orientation relationship that the c-austenite and d-ferrite share both in the base metal and in the fusion zone, despite the scarcity of c-austenite there. The c/d interphase boundaries that exhibit these well-known orientation relationships are highlighted in either red or blue (Figure 2, lower left). The crystallographic character of these interphase boundaries is typical of a Widmanstätten type transformation mentioned previously. Conversely, the high cooling rate (10 4 K( C)/s) in the fusion zone was sufficient to almost completely suppress all mechanisms of the solid-state transformation to c-austenite, resulting in an essentially singlephase d-ferrite fusion zone (Figure 2, lower right). A much different fusion zone microstructure is observed in alloy 10Mo13Cr12Ni under faster cooling conditions (42 mm/s weld, 10 5 K ( C)/s) (Figure 3) despite exhibiting the same Widmansta tten austenite morphology in the base metal. As shown in the EBSD phase identification map in Figure 3(a), the fusion zone is almost completely austenitic, even though the base metal (located at the bottom of each image) provides unequivocal evidence that this alloy exhibits primary ferrite solidification at lower cooling rates. The presence of a small amount of residual d-ferrite was observed at the fusion line (Figure 3(a)), so called because the d-ferrite shares the identical orientation as the d-ferrite in the base metal (Figure 3(b), orientation map), as would be expected from epitaxial growth at the fusion line during weld solidification. In contrast, the singlephase c-austenite exhibits no common orientation with 704 VOLUME 42A, MARCH 2011

6 Fig. 2 EBSD microstructural analysis of laser weld prepared at 4.2 mm/s on F mode alloy 10Mo18Cr17Ni. (Top) Light optical microscope image of fusion zone, (bottom left) orientation relationship map between c-austenite (band contrast, blue = N-W, red = K-S) and d-ferrite (white phase), and (bottom right) phase identification map indicating c-austenite (blue) and d-ferrite (red) regions of the fusion zone. the c-austenite in the base metal (Figure 3(c)), as one might expect if the fusion zone c-austenite nucleated and grew from the same d-ferrite grain or was the product of epitaxial solidification. While the relatively flat d/c interface observed in this alloy suggests the presence of an orientation relationship of some kind, analysis of the orientation of the c-austenite and d-ferrite across that boundary demonstrates the complete absence of a known orientation relationship (Figure 3(d)). The metallographically featureless appearance and absence of an orientation relationship between parent (d) and product (c) phases strongly suggests that the single-phase c-austenite observed in these F mode alloy welds is the result of a massive transformation. However, the hallmark of massive phenomena is the compositional invariance between product (c) and parent (d) phases resulting from the lack of long-range diffusion during the transformation. Consequently, EPMA was used to measure composition as a function of distance across both c-austenite and d-ferrite microstructural features. A step size of 2 lm was used in agreement with previous studies [22,23] that have reported the presence of massive c-austenite in Fe-Ni-Cr alloys at even higher cooling rates. While it is recognized that the cell spacing of the substructural features being measured is similar to the step size, it was concluded that long scans (>100 lm) should provide a sufficient statistical basis to detect significant compositional variation at boundaries if present. To establish a benchmark for evaluation of the large c-austenite grains, fully ferritic microstructures were Fig. 3 EBSD analysis of massive transformation microstructure observed in fusion zone of 42 mm/s weld prepared in F mode alloy 10Mo13Cr12Ni. (a) Phase ID map, (b) d-ferrite orientation map (similar colors = similar orientations), (c) c-austenite orientation map, (d) relative orientation map between c-austenite and d-ferrite. Note the lack of orientation relationship (OR) between massive c-austenite and the parent d-ferrite in the fusion zone. analyzed to confirm the ability of primary ferrite solidification to produce compositionally homogenous microstructures. The EPMA results for a 4.2 mm/s weld on 8Mo17Cr13Ni, displayed in Figure 4, indicate a uniform distribution of Mo, Ni, Cr, and Fe at their nominal levels. This is but one example of the analysis conducted on several other F mode laser weld microstructures prepared at multiple travel speeds, and confirms the diffusive ability of d-ferrite to eliminate concentration gradients that may occur during solidification, even at extremely high cooling rates (10 5 K( C)/s). When the single-phase austenitic structures such as those observed in a 4.2 mm/s laser weld prepared on 10Mo13Cr12Ni were examined (Figure 5), the uniform distribution of Mo, Cr, and Ni was demonstrated to be identical to that in residual d-ferrite near the weld fusion line. When compared to the highly variable chemical distribution in an arc button of a 6 wt pct Mo alloy exhibiting primary c-austenite solidification, presented previously by Anderson et al., [21] it is immediately obvious that the laser weld microstructure must be the product of a massive transformation from d-ferrite and not c-austenite solidification. The presence of Mo, Cr, and Ni in c-austenite at their nominal levels also eliminates long-range diffusion mechanisms from consideration. EPMA traces were collected for all F mode VOLUME 42A, MARCH

7 Fig. 4 EPMA data for a scan across fully ferrite microstructure of laser weld prepared at 4.2 mm/s on F mode alloy 8Mo17Cr13Ni. alloys exhibiting this type of single-phase c-austenite morphology, and the chemical distribution within c-austenite was consistently observed to be uniform. The reader will note that two F mode alloys, 0Mo20Cr10Ni and 8Mo13Cr12Ni, exhibited a minor shift in solidification mode from F to FA during laser welding. The close proximity of these compositions to the eutectic triangle [21] suggests that even small variations in composition could result in the formation of small amounts of eutectic c-austenite at the end of solidification, thereby changing the morphology of the massive c-austenite that forms. B. Ferrite-Austenite Solidification Mode Unlike the F solidification mode, ferrite-austenite (FA) solidification is characterized by primary d-ferrite cells with eutectic c-austenite forming at the cellular boundaries. The increasing relative thermodynamic stability of c-austenite on cooling causes the c-austenite present at the intercellular boundaries to consume the d-ferrite by epitaxial growth, thereby generating a room-temperature microstructure with much more c-austenite than is present at the end of solidification. The morphology of the residual d-ferrite, shown elsewhere in multiple sources, [12,21 23] is referred to as vermicular or skeletal due to the presence of a long spine of untransformed d-ferrite along dendrite cores with side branches or ribs extending into the adjacent c-austenite. Laser welds on FA alloys exhibited a refinement of the microstructure, as shown in Figure 6(a) for a weld on alloys 2Mo18Cr12Ni prepared at 4.2 mm/s. The microstructures shown are distinct in appearance, because d-ferrite is present at dendrite cores instead of the boundaries (as is the case for AF solidification). As shown in Figures 6(b) and (c), regions of the skeletal ferrite microstructure are interrupted by regions of a single-phase exhibiting a patchy appearance. Even at high magnifications, these featureless regions have poorly defined metallographic boundaries with the duplex d-ferrite + c-austenite structure, and small remnants of solidification substructure are observed within the single-phase region. The transformation from d-ferrite to c-austenite was observed to approach completion (i.e., the area fraction of regions 706 VOLUME 42A, MARCH 2011

8 Fig. 5 EPMA data for a scan across massive c-austenite and residual d-ferrite in laser weld prepared at 4.2 mm/s on F mode alloy 10Mo13Cr12Ni. exhibiting the patchy appearance increased) with increasing cooling rate, as indicated by the behavior exhibited by alloy 4Mo20Cr15Ni in Figure 7, and was similar to the behavior observed with the F mode alloys discussed previously. The lathy ferrite morphology is also observed in several welds near the fusion line (Figure 7), distinct from skeletal ferrite because of the faceted appearance of the d/c phase boundary caused by the K-S orientation relationship between austenite and ferrite in the lathy morphology. As demonstrated by Brooks et al. [33] and Inoue et al., [34] c-austenite nucleates on d-ferrite during solidification in this mode, thereby establishing the K-S relationship, which is maintained during growth. This orientation relationship results in faceting at the c/d interface. All alloys observed to solidify in the FA mode under arc button melting conditions exhibited this type of single-phase c-austenite to some extent at the high cooling rates associated with laser welds (Table V). Phase identification via EBSD analysis positively determined this single phase to be c-austenite, an example of which is displayed in Figure 8 for a 42 mm/s weld prepared on alloy 10Mo14Cr16Ni. The accompanying austenite orientation map demonstrates the large patchy appearance of the c-austenite grains, along with a large variation in relative orientation between grains in the fusion zone and the base metal. Thus, as shown for F mode alloys previously, the singlephase austenite fusion zone is not the product of epitaxial growth from the base metal during weld solidification. EPMA compositional data were also collected on all FA mode laser welds exhibiting this type of microstructure, an example of which is shown in Figure 9 for a 21 mm/s weld on alloy 10Mo14Cr16Ni. As seen for massive c-austenite in the three F mode alloys previously, uniform distributions of Mo, Cr, Ni, and Fe are observed within the austenitic structure, at their nominal levels. VOLUME 42A, MARCH

9 Fig. 6 Examples of massive c-austenite morphology for laser weld prepared at 4.2 mm/s on FA mode alloy 2Mo18Cr12Ni: (a) refined vermicular ferrite microstructure at center of fusion zone, (b) and (c) c m (light phase) growth in the fusion zone, and (c) immediate juxtaposition of massive c and vermicular d morphology. C. High Resolution Compositional Mapping As mentioned previously, the EPMA data collected in this study on laser welds in both F mode and FA mode alloys were the product of line scans with a step size of 2 lm and an interaction volume of approximately 1 lm 3. The choice of step size allowed for the analysis of a significant amount of material in a reasonable amount of time, but a source of concern is that the interaction volume of this analytical technique is similar in size to the experimentally measured dendrite spacing (1.5 to 3 lm) in the laser welds. Specifically, this could potentially result in artificial averaging of the composition distribution, leading to the errant conclusion that the single-phase c-austenite has the same uniform distribution as the parent d-ferrite. To address this issue directly, three alloys with similar compositions (10Mo14Cr21Ni, 10Mo14Cr16- Ni, and 10Mo13Cr12Ni) but different primary solidification modes (AF, FA, and F, respectively) were subjected to compositional mapping using a transmission electron microscope with nanometer resolution. Thin film specimens were prepared from the 42 mm/s (10 5 K ( C)/s) welds on each alloy using the FIB milling process. The highest cooling rate was specifically chosen as a worst-case scenario in that the time allowed for solute diffusion on cooling would be minimized and (arguably) the potential for partitionless solidification, which might be suggested as an alternative explanation for the microstructures observed here, would be maximized. The resulting maps taken with a step size of 9 nm are shown in Figure 10. The color scheme in each image is based on the relative intensity of the elements in the mean spectrum for the area being analyzed at each point/pixel. The numbers shown for each legend give the reader a feel for the relative differences between the color maxima and minima; i.e., while the same absolute difference exists between 0.5 and 1 on a 0 1 scale and between 1.5 and 2.0 on a 1 2 scale, the relative intensities of each spectrum are much more pronounced on the 0 1 scale. In other words, changing the scale from 0 1 to higher values amplifies small relative differences that would not otherwise be noted on a 0 1 contrast scale. Thus, in Figure 10(a), the primary austenite dendrites of AF mode alloy 10Mo14Cr21Ni are shown to be depleted in Mo (blue) relative to the interdendritic boundaries (red), caused by microsegregation during solidification. Figures 10(b) and (c) compare the chemical distribution of d-ferrite and c-austenite, respectively, for F mode alloy 10Mo13Cr12Ni. The only source of contrast seen in these maps was differences in thin film thickness, indicated by the gradation observed across the sample. Similarly, the austenite from FA alloy 10Mo14Cr16Ni (Figure 10(d)) exhibits only thickness contrast and has a uniform distribution of Mo, Cr, and Ni, further confirming the massive nature of the austenite in the laser weld fusion zone of these alloys. Note that this thickness contrast was only visible because of the shift in contrast scale above 0 1. IV. MASSIVE TRANSFORMATION AFTER FA SOLIDIFICATION The unique situation presented by the FA solidification mode is the termination of solidification with c-austenite present at d-ferrite cell boundaries. Nucleation of c-austenite is no longer required to achieve the decomposition of d-ferrite on cooling as in the case of F mode alloys. However, the absence of a barrier to nucleation makes the suppression of competing 708 VOLUME 42A, MARCH 2011

10 Fig. 7 Transformation to c-austenite observed to go further toward completion with increasing cooling rate and solidification velocity in FA alloy 4Mo20Cr15Ni: (a) and (b) 4.2 mm/s, (c) and (d) 21 mm/s, and (e) and (f) 42 mm/s. Residual lathy d-ferrite structure observed at fusion line at 21 and 42 mm/s. long-range diffusion mechanisms exceedingly difficult, but maybe not impossible. Two questions must therefore be answered: (1) Can a massive transformation take place without nucleation, with a quench below T 0, to induce enough undercooling such that massive c can grow epitaxially from the pre-existing austenite? (2) Could the observed single-phase c-austenite be the result of some other phenomenon such as partitionless solidification? A. Massive Growth without Nucleation Because massive transformations are diffusional in nature, interfacial mobility is maximized through incoherency of the advancing transformation boundary; a distinguishing characteristic of the massive transformation is, therefore, the lack of an orientation relationship between parent and product. [1,35 37] The high mobility of these interfaces and the unique ability of the massive transformation to cross parent grain boundaries unimpeded [38] is consistent with the absence of an orientation relationship. [1,39] Even though massive phase boundaries often exhibit planar facets, which suggest an orientation relationship between parent and massive product, none has been reported. Considerable controversy exists regarding the true nature of the massive/parent interface and the level of coherency that may be present, [3,5,7,40 42] and none of the data presented in this study are intended to address these complex mechanistic issues. However, it should be noted that VOLUME 42A, MARCH

11 Fig. 8 Phase identification of massive c-austenite (left) and c-austenite orientation map (right) in a laser weld prepared at 4.2 mm/s on FA alloy 10Mo14Cr16Ni. these mechanistic studies focused on binary and ternary alloy systems, many of which have limited applicability to the Fe-Ni-Cr-Mo studied here. Inoue et al. [43] claimed that massive c maintains a K-S orientation relationship with parent d-ferrite, but the nature of c-austenite studied was inconsistent with traditional observations about massive phenomena. In particular, distinct differences in composition were observed between parent d-ferrite and product c-austenite, suggesting that partitioning occurred during the transformation hence, no massive transformation. Several authors have postulated that massive growth can occur without nucleation (i.e., from a previously existing region of the product phase). Plichta [44] suggested that it may be reasonable to consider massive growth to be possible from the fastest growing nuclei. Fig. 9 EPMA data for a scan across massive c-austenite in laser weld prepared at 21 mm/s on FA mode alloy 10Mo14Cr16Ni. 710 VOLUME 42A, MARCH 2011

12 Fig. 10 AEM maps of (a) AF mode alloy 10Mo14Cr21Ni, (b) d-ferrite in F mode alloy 10Mo13Cr12Ni, (c) c-austenite in F mode alloy 10Mo13Cr12Ni, and (d) c-austenite in FA mode alloy 10Mo14Cr16Ni collected with a step size of 9 nm. The gradation observed in (b) through (d) is exclusively the result of thickness contrast. He suggested that there is no difference between a massive or equilibrium nucleus; growth can still occur massively, and the product phase is only determined by growth kinetics. The best experimental example in the literature of this may be the Cu-Ga-Ge system, [45,46] which exhibits a narrow two-phase region where the b (bcc) and f (hcp) phases coexist. At high temperature, the microstructure typically consists of b regions outlining f grains. Upon rapid quenching of such a structure, the hexagonal f grains remain stable, but each b region transforms via a massive transformation to an hcp f m structure without change in composition. A coherent f/f m boundary was observed [45] that exhibited a lattice parameter discontinuity across it due to the difference in composition between the previously existing f and the f m that grows from it at the b composition. This growth without nucleation feature demonstrates that massive transformations exhibit both difficult nucleation and easy growth. [46] The challenge is establishing parent/ product interface with a sufficiently incoherent character to allow easy growth, and this is why the massive kinetics are so dependant on the rate of nucleation. If such interfaces are already present, they will be used without the need for nucleation. Similar behavior was also reported in the Cu-Al binary system. [47] The present experimental data indicate that the singlephase c-austenite observed in the fusion zone of each FA alloy was the result of a massive transformation without nucleation and not the completion of long-range diffusional growth. The transformation from d-ferrite to c-austenite was observed to systematically approach completion more readily as the cooling rate increased, whereas just the opposite would be expected for diffusional growth. Less time is available for diffusion at high cooling rates, despite the fact nucleation is not required and that the required diffusion distance decreases with the reduction of the microstructural scale. Transformations involving long-range diffusion also exhibit compositional differences caused by partitioning between parent and product phases. The uniform chemical distributions exhibited by the EPMA and AEM data collected for these alloys demonstrate that such partitioning was absent. Finally, a moving massive c-austenite interface will be arrested when it impinges on other c-austenite grains (depicted schematically in Figure 11), which explains the patchy appearance of the single-phase austenite in several alloys. Thus, based on the experimental data presented here, the single-phase c-austenite observed in FA mode alloys is only consistent with massive growth without nucleation. VOLUME 42A, MARCH

13 Fig. 12 Calculated vertical phase diagram with constant additions of 6 wt pct Mo and 64 wt pct Fe used to demonstrate thermodynamic arguments for kinetic shift in primary solidification mode and partitionless c-austenite solidification. Top red line = metastable c liquidus (schematic), bottom red line = metastable c solidus (schematic), and dashed green line = T 0 for liquid + c two-phase region (schematic). a common temperature gradient. [34] The incoherent d/c interface in skeletal d-ferrite is inherently more mobile than the coherent interface in lathy d-ferrite and can be induced to rapid massive growth without nucleation. This explains the well-defined and repeatable nonmassive region in FA mode alloys at the fusion line. Fig. 11 Schematic representation of the origin of microstructures in FA mode alloys. (a) Dendritic d-ferrite forms during FA solidification with intercellular c-austenite, (b) skeletal d-ferrite morphology results from diffusive growth of intercellular c-austenite into d-ferrite, and (c) epitaxial growth of massive c-austenite without nucleation. The thin layer of residual FA mode microstructure consistently observed at the root of FA laser welds exhibiting massive c-austenite (Figure 7) can be attributed to the morphology of d-ferrite in this regime. As shown by Inoue and co-workers, [48] the beginning of the FA solidification mode is frequently characterized by the nucleation of d-ferrite on the small amount of c-austenite that often forms epitaxially at the fusion line in the initial stages of solidification. At this juncture, the K-S orientation relationship is established and is maintained at the early stages of solidification, resulting in the formation of lathy ferrite near the fusion line. As the relative orientation of the maximum temperature gradient continues to change, it becomes increasingly difficult for d-ferrite to align its easy growth direction with this gradient while simultaneously maintaining a planar orientation relationship with c-austenite. Consequently, a shift to the skeletal ferrite morphology is quickly favored at the high cooling rates and solidification velocities associated with the laser welds prepared in this study because the c/d interface in this morphology is incoherent. However, the easy growth directions of primary d-ferrite and intercellular c-austenite will naturally be parallel because they will both be aligned along B. Partitionless Solidification Identification of massive structures is often based on circumstantial evidence: composition invariance between parent and product, rapid velocity of the transformation interface, and evidence of an incoherent interface between parent and product. Because of this, past work by Vitek and David [49] suggested that partitionless solidification could also produce similar single-phase structures. This postulation has a thermodynamic basis for processes with high cooling rates, such as HED welding, because partitionless solidification depends on much the same thermodynamic criteria as does a shift in solidification mode from primary d-ferrite to primary c-austenite. This can be explained schematically by viewing the phase diagram based on an alloy with 6 wt pct Mo and a Ni + Cr content of 30 wt pct shown in Figure 12. Shifts to primary c-austenite solidification are thermodynamically rationalized by the argument that high cooling rates in near-eutectic alloys (C 01 ) can induce enough undercooling to suppress primary d-ferrite solidification until the temperature reaches the extended metastable c-austenite liquidus (top red line in Figure 12). A similar argument is made for partitionless solidification, except that undercooling to the extended liquidus is inadequate; rather, a minimum undercooling to the T 0 temperature (green dashed line, Figure 12) in the liquid + c-austenite two-phase region is necessary. In many ways, this is an analog to the massive transformation in the solid state between d-ferrite and c-austenite, except that massivelike c-austenite forms directly from the liquid. 712 VOLUME 42A, MARCH 2011

14 While Brooks and co-workers [11,12] have acknowledged the technical merit of these thermodynamic arguments for kinetic shifts in solidification mode, they have dismissed partitionless solidification as an explanation for the single-phase c-austenite structures observed in HED welds in the Fe-Ni-Cr system. The same conclusions are reached here as well. Partitionless solidification requires the planar growth of c-austenite, which would be unimpeded by the amorphous structure of the liquid phase. However, planar c-austenite was only periodically observed at the fusion line of the welds in previous studies in this material system, [25] indicative of the early stages of epitaxial solidification before the onset of constitutional supercooling that generates the dendritic substructure seen elsewhere in the fusion zone. Planar growth was limited only to this regime, and the morphology of the massive c-austenite boundaries is far from being completely planar. Furthermore, calculations made in similar systems [50 52] demonstrated that the switch in primary solidification mode from d-ferrite to c-austenite occurs at high velocities and the transition to planar c-austenite at the absolute limit of stability requires even higher levels. These values are well in excess of the conditions studied here. Even at travel speeds of 5000 mm/s and cooling rates more than an order of magnitude higher than those observed here (10 6 K ( C)/s), partitionless solidification was not observed in the Fe-Ni-Cr system. [22] Moreover, it seems unlikely that partitionless solidification would be responsible for the patchy appearance observed on such a small level in FA modes, where partitionless c-austenite is immediately adjacent to regions of vermicular ferrite. This also does not consider the intermittent crossing of d-ferrite boundaries in F mode alloys (Figure 1(b)) and the massive c-austenite grain orientation that is inconsistent with planar growth. It should be noted that some of the kinetic arguments that are used to describe partitionless solidification were invoked by Lima and Kurz [6] to describe the solid-state massive transformation mechanism, but the present study is not intended to clarify or address the issues raised there. From a thermodynamic standpoint, partitionless solidification can also be eliminated from consideration based on experimentally observed trends. Examination of Figure 12 reveals that the undercooling required to induce partitionless c-austenite solidification in primary d-ferrite alloys is very large, as quenching past the T 0 temperature in the metastable c-austenite + liquid two-phase region (C 01, C 02 ) is required. Such large undercoolings were not present under the conditions examined in this study, because no shift in primary solidification mode was observed in the alloys discussed here, despite the smaller undercooling required to achieve such a shift in alloys near the eutectic (C 01 ). Moreover, partitionless c solidification would only be possible for primary d-ferrite alloys if it were also observed in primary c-austenite, where the undercooling required to switch to partitionless solidification is much less. A parallel study [21] reported no evidence of this behavior reinforcing the conclusions reached here that the single-phase c-austenite structures can only be attributed to massive transformation phenomena. V. THERMODYNAMICS AND MASSIVE TRANSFORMATION It is well known that massive transformation can only occur once an alloy is quenched below the critical temperature, T 0. [1 3,6,10] This value is defined as the temperature for a given composition where the free energy of one phase in a two-phase region is equal to the other. For the Mo-bearing stainless steels in this study, c-austenite achieves a lower free energy than d-ferrite once the temperature drops below T 0 in the c-austenite + d-ferrite two-phase field. The thermodynamics of the massive transformation are satisfied once this occurs, but like other diffusional processes, the driving force for massive nucleation and growth increases significantly as undercooling is increased below T 0.In fact, it was often surmised that quenching into the single-phase c-austenite region was required for the massive transformation to occur. However, work by Massalski [1] and Hillert [2,53] demonstrates this is not a requirement. In fact, Menon et al. [54] suggest that massive transformation in a two-phase field below T 0 is probably viable. This particularly pertains to the case when volume diffusivity in the matrix required for appreciable solute partitioning during growth is significantly lower than the trans-interphase boundary diffusivity required for the massive transformation. The sluggish substitutional diffusivity of Mo in c-austenite [55,56] would certainly seem to qualify. Massive transformations are expected to be favored at higher absolute values of T 0, where higher undercoolings can be achieved in a temperature range where the mobility of solute atoms is appreciable. Elmer [22,23] reported a minimum threshold travel velocity for the massive transformation at 100 mm/s in the Fe-Ni-Cr system, corresponding to a cooling rate of K( C)/s, but only observed massive c-austenite in one particular alloy (out of seven). The required undercooling was shown to be extremely dependent on the nominal composition as the T 0 line was found to have a steep slope with composition ( 85 K ( C)/wt pct Cr). Analysis of the kinetics of the massive transformation in Fe-Ni- Cr alloys was also presented, where the diffusivity of Ni in c (slower than Fe and Cr) was shown to be the ratecontrolling variable. This conclusion implies that in Mo-bearing stainless steels, the low diffusivity of Mo in c should significantly affect the kinetics of the massive transformation. Based on this information, both the absolute temperature of T 0 and the undercooling (DT 0 ) in the two-phase region required to reach T 0 while suppressing competing diffusional processes would seem to be important criteria for inducing a massive transformation. Thus, the required thermodynamic values were calculated for each alloy in this study using Thermo-Calc [57] thermodynamic software coupled with the Fe-Data thermodynamic database. [58] Experimentally measured average compositions were used, resulting in calculated values of liquidus (T m ), terminal solidus, T 0, and DT 0. A vertical section of the Fe-Ni-Cr-Mo phase diagram with 2 wt pct Mo, 68 wt pct Fe is displayed in Figure 13 to schematically illustrate the variability of T 0 and DT 0 as a VOLUME 42A, MARCH