Effects of Carbon Content and Austenite Grain Size on Retained Austenite Fraction in Stir Zone of Friction Stir Welded 6%Ni Carbon Steels

Abstract

Friction stir welding (FSW) was performed under the two welding conditions (rotation speed-traveling speed) of 150 rpm–100 mm/min and 200 rpm–400 mm/min using 6 mass%Ni steels with different carbon contents from 0.14 mass% to 0.63 mass%. The slightly lower peak welding temperature and the higher cooling rate were predicted under the condition of 200 rpm–400 mm/min. The effects of carbon content and prior austenite grain size on retained austenite fraction in the stir zones were evaluated. When carbon content was 0.30 mass% or more, a fine microstructure consisting of lath martensite and retained austenite was formed in the stir zone. Irrespective of welding conditions, the amount of retained austenite increased with the increase of carbon content. More retained austenite was obtained at the stir zone under the 200 rpm–400 mm/min condition, which resulted from rapid cooling and finer prior austenite grains compared with the condition of 150 rpm–100 mm/min. The effect of the prior austenite grain size on the amount of retained austenite was successfully extracted, and the finer austenite grain size was concluded to play an important role to stabilize austenite by analyzing the difference in the martensitic start temperatures predicted based on the chemical composition and the retained austenite fraction based on Koistinen-Marburger equation.

1. Introduction

Friction stir welding (FSW) is a solid-state welding technique that utilizes the frictional heat and stirring force generated by a high-speed rotating cylindrical tool pushing against joining members. Applications of low-melting-point metals such as Al alloys were investigated^1,2,3,4,5) when FSW was first introduced. Recently, applications for high-melting-point metals such as steels have been investigated owing to improvements in tool materials.^{5,6,7,8,9,10,11,12)} In many cases involving the FSW of steels, a phase transformation occurs via frictional heating and cooling during and after FSW. In other words, austenite which transforms from ferrite during stirring undergoes dynamic recrystallization, followed by phase transformation from austenite to ferrite and/or martensite during cooling after the stirring stage. Because the phase transformation behavior after stirring is affected by the thermomechanical history, such as the plastic deformation, maximum temperature, and cooling rate, microstructures in the stir zone vary based on the phase transformation behavior, which affects the mechanical properties of the joints. Therefore, it is important to clarify the microstructure formed during and after stirring and understand the factors affecting the phase transformation behavior during cooling. However, evaluating the microstructure of austenite before phase transformation is difficult because the microstructure changes due to phase transformation. Consequently, many aspects have remained unclarified.

The authors reported that austenite in the stir zone was stabilized and retained after FSW under appropriate welding conditions, which resulted in an improvement in the strength-elongation balance of the stir zone owing to the transformation-induced plasticity (TRIP) effect in low-alloyed^12,13) and alloyed steels.^14,15,16) Although austenite stability is assumed to be affected by the dislocation density and grain size of the prior austenite and the mobility of the phase transformation interface,¹⁷⁾ the mechanisms affecting austenite stability in terms of concrete welding conditions and fraction of retained austenite have not been elucidated.

Therefore, in this study, FSW was performed under two welding conditions using 6 mass%Ni steels with different carbon contents from 0.14 to 0.63 mass%, and the amount of retained austenite at room temperature was investigated. Subsequently, the effects of carbon content and austenite grain size on the retained austenite fraction in the stir zone were discussed based on the welding conditions.

2. Materials and Methods

2.1. Materials

Ingots prepared by adding Ni and graphite to the starting material of a low-carbon cold-rolled steel using a high-frequency induction melting furnace were formed into sheets with a thickness of 1.6 mm via hot rolling at approximately 950°C, followed by air-cooling to room temperature. The chemical compositions of the sheets with different carbon contents were evaluated using an optical emission spectrometer (SHIMADZU, PDA-7000), as shown in Table 1. In this study, each material was referred to as 6%Ni-x_C%C (mass%) steel based on the amount of Ni and carbon x_C. Figure 1 shows the 500°C and 730°C isothermal sections of the Fe–C–Ni ternary phase diagram.¹⁸⁾ Irrespective of the materials, the structure comprising ferrite and (Fe, Ni)₃C was stable at low temperatures (Fig. 1(b)), whereas at 730°C, which is approximately the welding temperature for the present study, the austenitic single phase was stable, except for 6%Ni–0.14%C steel (Fig. 1(a)).

Table 1. Chemical compositions of 6%Ni-x_C%C steels used (mass%).

Alloy	C	Si	Mn	Ni	Cr	Fe
6%Ni-0.14%C	0.14	0.03	0.16	5.75	0.02	bal.
6%Ni-0.30%C	0.30	0.03	0.17	5.86	0.02	bal.
6%Ni-0.46%C	0.46	0.02	0.16	5.85	0.02	bal.
6%Ni-0.54%C	0.54	0.02	0.16	5.76	0.02	bal.
6%Ni-0.56%C	0.56	0.03	0.16	5.97	0.02	bal.
6%Ni-0.63%C	0.63	0.02	0.17	5.82	0.02	bal.

Fig. 1.

Isothermal sections of Fe–C–Ni phase diagrams¹⁸⁾ at (a) 730°C and (b) 500°C. The compositions of present study are indicated by rhombus marks.

2.2. Welding Conditions

FSW was performed on a single sheet (stir-in-plate) using a tool made of cemented tungsten carbide with a shoulder diameter, probe diameter, and probe length of 12, 4, and 1.4 mm, respectively. Two welding conditions involving a tool rotation speed and traveling speed of 150 rpm–100 mm/min (15.7 rad/s–1.67 mm/s) and 200 rpm–400 mm/min (20.9 rad/s–6.67 mm/s), respectively, were employed for the FSW. These two welding conditions were regarded as the low and high cooling rate conditions, respectively, with the maximum temperature of austenite single-phase. In our previous studies,^8,19) the maximum temperatures under the conditions of 150 rpm–100 mm/min and 200 rpm–400 mm/min were predicted to be 730°C and 700°C, whereas the cooling rates were predicted to be 60°C/s and 100°C/s, respectively. The bead length was 40 mm, and the tool plunging load was maintained at 25.5 and 36.3 kN under the conditions of 150 rpm–100 mm/min and 200 rpm–400 mm/min, respectively. Herein, the sheet normal, welding, and transverse directions are denoted as ND, WD, and TD, respectively. Furthermore, the side where the tool rotating and traveling directions are the same is defined as the advancing side (AS), whereas the side where they are opposite is defined as the retreating side (RS).

2.3. Evaluation Methods

In the cross section of the stir zones perpendicular to the WD, macrostructure observation was performed via optical microscopy. Furthermore, microstructure observation via electron backscattering diffraction (EBSD) measurement was performed using scanning electron microscopy (SEM). Specimens were first prepared via electric discharge wire cutting, followed by mechanical polishing; subsequently, they were employed for optical microscopy observation after being subjected to chemical etching with 3% nital solution. Additionally, the specimens were employed for EBSD measurements after electrical polishing at 20 V using a HClO:CH₃COOH = 1:9 solution. EBSD measurements were performed at an acceleration voltage of 15 kV and a scanning step size of 0.05 μm in a 20 μm × 40 μm area located at a depth 0.7 mm from the top surface of the welding center. Measurement points with confident index (CI) values lower than 0.05 were excluded and are shown as black points in the maps. Moreover, using a micro-area X-ray diffractometer (D8-DISCOVER with GODDS, Bruker), diffraction profiles were obtained in the same area as the EBSD measurement in the cross section perpendicular to the WD using the Co–Kα radiation (λ = 1.79026 nm) generated at an acceleration voltage of 35 kV and a collimator diameter of 500 μm. The carbon content in austenite x_{C in γ} was calculated by substituting the lattice parameter of austenite a obtained from the measured diffraction profiles into the equation, a = 357.3 + 3.3 x_{C in γ} (mass%).²⁰⁾

3. Experimental Results

3.1. Microstructure of Base Metals

Figure 2 shows the SEM images of the cross-sectional microstructures of the base metals with carbon contents from 0.14 to 0.63 mass%, whereas Fig. 3 shows the phase distribution maps of the base metals with carbon contents of 0.30 and 0.63 mass%. The body-centered cubic (BCC) phase comprising ferrite and martensite, as well as the face-centered cubic (FCC) phase comprising austenite are shown as dark gray and light gray areas in Fig. 3, respectively. Here, high-angle boundaries with a misorientation angle exceeding 15° are indicated by a black line. As shown in Fig. 2, each material was composed of a mixed microstructure comprising equiaxed ferrite, pearlite, martensite, and retained austenite, and the proportion of martensite increased with the carbon content. Furthermore, the retained austenite distributed primarily in the martensite, as shown in Fig. 3.

Fig. 2.

Secondary electron images of base metals of 6%Ni- (a) 0.14%C, (b) 0.30%C, (c) 0.46%C, and (d) 6%Ni-0.63%C steels.

Fig. 3.

Phase maps of base metals for 6%Ni- (a) 0.30%C, (b) 0.63%C steels. The distributions of BCC and FCC phases. The area fractions of these phases are denoted on each image.

3.2. Macrostructure and Hardness Distribution around Stir Zones

Figure 4 shows optical macrostructural images in cross sections perpendicular to the WD of 6%Ni–0.30%C steel, 6%Ni–0.46%C, and 6%Ni–0.63%C steel friction stir welded at each welding condition. All images are organized such that the AS is located on the right side. Welding defects were not observed on the weld surface and in the cross-sectional images of all the joints, as shown in Fig. 4. Areas exhibiting contrast levels different from that of the base metal were identified as the stir zones, which suggests the occurrence of phase transformation during welding. Figure 5 shows the Vickers hardness distributions along the TD in the cross sections of 6%Ni–0.30%C and 6%Ni–0.63%C steels friction stir welded at each welding condition. The hardness in the weld center region of the stir zone, where phase transformation occurred, was higher than those of the base metal and heat-affected zone, and its distribution was almost uniform. By contrast, the hardness decreased in the heat-affected zones located outside the stir zone of all welds.

Fig. 4.

Cross-sectional images of (a, d) 6%Ni-0.30%C, (b, e) 6%Ni-0.46%C, and (c, f) 6%Ni-0.63%C steels FSWed at (a–c) 150 rpm–100 mm/min and (d–f) 200 rpm–400 mm/min.

Fig. 5.

Distributions of Vickers hardness along TD in cross sections of 6%Ni-0.30%C and 6%Ni-0.63%C steels FSWed at 150 rpm–100 mm/min and 200 rpm–400 mm/min.

3.3. Microstructure and Retained Austenite in Stir Zones

Figure 6 shows the SEM images of stir zones of materials composed of different carbon contents friction stir welded under the condition of 150 rpm–100 mm/min. Figure 6(a) shows the mixed microstructure comprising ferrite, pearlite, and martensite with diameters of a few micrometers in the stir zone of 6%Ni–0.14%C. In the materials with carbon content exceeding 0.30 mass%, most regions contained martensite and retained austenite. The crystal orientation maps of the BCC and FCC phases obtained via SEM–EBSD in the center of the stir zones are shown in Fig. 7. Crystal orientations parallel to the WD of the BCC and FCC phases are shown separately. The microstructure of 6%Ni–0.14%C steel (Figs. 7(a) and 7(e)) was composed of a mixture of equiaxed ferrite grains with low misorientation and an elongated martensite, which corresponds to the SEM image in Fig. 6(a). Moreover, very small amount of fine granular retained austenite was observed between BCC grains. Steels with higher carbon contents exceeding 0.30 mass% exhibited microstructures composed primarily of martensite with elongated shapes, in which retained austenite was distributed. It is noteworthy that the characteristics of the microstructures formed during welding differed significantly between 6%Ni–0.14%C steel and other steels with higher carbon contents. Moreover, the microstructures of the stir zones friction stir welded under the condition of 200 rpm–400 mm/min, which were subjected to a higher cooling rate after FSW, comprised finer martensite structures with a greater amount of retained austenite compared with the microstructures of the stir zones friction stir welded under the condition of 150 rpm–100 mm/min. In addition, irrespective of the welding condition, the area fraction of retained austenite increased with the carbon content of the base metals.

Fig. 6.

Secondary electron images of stir zones FSWed at 150 rpm–100 mm/min of (a) 6%Ni-0.14%C, (b) 6%Ni-0.30%C, (c) 6%Ni-0.46%C, and (d) 6%Ni-0.63%C steels.

Fig. 7.

Orientation color maps the center in of stir zones FSWed at (a–d) 150 rpm–100 mm/min and (e–h) 200 rpm–400 mm/min for 6%Ni- (a, e) 0.14%C, (b, f) 0.30%C, (c, g) 0.46%C, and (d, h) 0.63%C steels. The crystal orientations parallel to WD are displayed for BCC and FCC phases; the area fractions of these phases are denoted on each image. (Online version in color.)

Figure 8 shows the retained austenite fraction obtained via EBSD in the stir zones. As shown in Fig. 7, the retained austenite area fraction increased under the welding condition of 200 rpm–400 mm/min, where the welding temperature was lower and the cooling rate was higher compared with the welding condition of 150 rpm–100 mm/min. Moreover, the increase in the retained austenite area fraction with the carbon content of the base metals was confirmed.

Fig. 8.

Relationship between carbon content of base metal and retained austenite fraction obtained by EBSD in stir zones of 6%Ni-x_C%C steels FSWed at 150 rpm–100 mm/min and 200 rpm–400 mm/min.

The carbon content in the retained austenite, evaluated via X-ray diffraction, is shown in Fig. 9 as a function of the base metal carbon content. In the base metals with a carbon content exceeding 0.30 mass%, the carbon content in the retained austenite was almost similar to that of the base metal, except for the case of 0.30 mass%C and 150 rpm–100 mm/min. This suggests that the carbon content became uniform due to stirring, which occurred at the temperature of the austenite single-phase, and that carbon redistribution did not occur during rapid cooling after the stirring. By contrast, in 6%Ni–0.14%C steel, the carbon content in the retained austenite increased significantly compared with that of the base metal, which suggests that carbon diffused and separated significantly between the α and γ phases. Therefore, it was presumed that 6%Ni–0.14%C steel underwent stirring in the inter-critical temperature region, and that diffusional transformation occurred during cooling after stirring. This assumption is consistent with the results of microstructural observations presented in Figs. 6 and 7. Hence, the equiaxed ferrite and pearlite observed in the stir zones of 6%Ni–0.14%C steel friction stir welded under the two welding conditions was presumably formed during stirring at the inter-critical temperature region followed by cooling, where the diffusion of carbon can occur.

Fig. 9.

Relationship between carbon content of base metals and the retained austenite obtained by XRD in stir zones of 6%Ni-x_C%C steels FSWed at 150 rpm–100 mm/min and 200 rpm–400 mm/min.

Moreover, in the case of the 6%Ni–0.30%C steel friction stir welded under the condition of 150 rpm–100 mm/min, it can be assumed that equiaxed ferrite formation and carbon partitioning to the γ phase occurred via partial diffusional transformation during cooling after stirring owing to the relatively low cooling rate. By contrast, the microstructures observed in most of the steels with carbon contents exceeding 0.30 mass% were composed primarily of lath martensite, in which retained austenite was distributed between the elongated martensite blocks, where diffusionless transformation was presumed to occur and the diffusion of carbon was expected to be limited. In the following section, we analyze steels with carbon contents exceeding 0.30 mass%, and discuss the effects of the microstructure of the prior austenite on the phase transformation during cooling after stirring in the austenite single-phase.

4. Discussion

4.1. Evaluation of Prior Austenite Grain Size in Stir Zones

The Kurdjumov–Sachs (K–S) orientation relationship holds between the parent phase of austenite and the product phase of martensite. As shown in Fig. 10, the K–S orientation relationship was confirmed between the retained austenite and martensite in the stir zones of 6%Ni–0.63%C steel. The prior austenite grain structure was reconstructed from martensite using the K–S orientation relationship.^21,22) The relationship between the carbon content of the base metal and the average grain intercept of the reconstructed prior austenite grains in the stir zones of 6%Ni–0.46%C and 6%Ni–0.63%C steels is shown in Fig. 11. The error bars indicate the standard deviation. As observed, finer austenite grains were formed under the condition of 200 rpm–400 mm/min, where the cooling rate was higher. In addition, although the effect of carbon content on the prior austenite grain size was less prominent compared with that of the welding condition, the prior austenite grain size in the stir zones increased with the carbon content, regardless of the welding condition. Although the reconstruction of a prior austenite structure is the best approach for evaluating the austenite grain size, a significant amount of time is required to reconstruct it by verifying the orientations of the retained austenite and every neighboring blocks. Moreover, reconstructing the entire region for analysis in cases involving extremely small prior austenite grains or a few retained austenite grains is difficult. Therefore, the prior austenite grain size of all the samples could not be evaluated. Hence, in this study, another method to evaluate the prior austenite grain size, which is simpler than the reconstruction of prior austenite structures, was employed.

Fig. 10.

An example of reconstruction of a prior austenite grain structure exploiting the orientation relationship between martensite and retained austenite. 001 pole figures of martensite and retained austenite contained in the prior austenite grain are shown.

Fig. 11.

Relationship between carbon content of base metal and average grain intercept of reconstructed prior austenite grain in stir zones of 6%Ni-0.46%C and 6%Ni-0.63%C steels.

In EBSD measurements, high-angle boundaries are evaluated more accurately than the low-angle boundaries. In the lath martensite structure, a block is the smallest unit and can be regarded as a crystal grain surrounded by relatively high-angle boundaries. Therefore, not only the block in the lath martensite structure but also the relationship between the block size and prior austenite grain size or the area fraction of retained austenite were investigated.

In the lath martensite structure, the boundaries between blocks exhibits a relatively high-angle misorientation of approximately 10°. By contrast, the misorientation of lath boundaries is only 1°–2°.^23,24) Hence, in the present analysis, all boundaries with misorientation angles exceeding 5° were considered to determine the block size. The arithmetic average of the block length, which is defined as the maximum feret length of each block as shown in Fig. 12, was evaluated in the stir zones of each sample and was plotted as a function of the carbon content of the base metals. In both welding conditions, the average block length decreased significantly as the carbon content increased from 0.3 to 0.46 mass%, and then increased slightly as the carbon content increased further from 0.46 to 0.63 mass%. This tendency is discussed in the following section. Regarding the effect of the welding condition on the block length, the average block length welded under the condition of 200 rpm–400 mm/min was smaller than that welded under the condition of 150 rpm–100 mm/min for all specimens with different carbon contents.

Fig. 12.

Relationship between carbon content of base metal and average block length in stir zones of 6%Ni-x_C%C steels.

Figure 13 shows the relationship between the average grain intercept of the prior austenite grains (Fig. 11) and the average block length (Fig. 12). The prior austenite grain size and average block length were revealed to exhibit a linear relationship. As such, the average block length decreased with decreasing the prior austenite grain size, regardless of the base metal carbon content and welding condition. Maki et al. indicated that packet size and block width decreased linearly with decreasing the prior austenite grain size in 0.2%C steel and 18%Ni steel.²⁵⁾ Based on a packet composed of parallel blocks, it was inferred that the tendency of the block length shown in this study reflects the variation in the packet size. In other words, it can be assumed that the block length reflects the prior austenite grain size. Hence, the tendency of block length shown in Fig. 12 indicates that the prior austenite grain size welded under the condition of 200 rpm–400 mm/min was finer than that welded under the condition of 150 rpm–100 mm/min for all the samples with different carbon contents. Furthermore, this tendency of the prior austenite grain size presumably reflects the variations in grain refinement and grain growth with the welding condition, owing to the difference in the stirring temperature and cooling rate.

Fig. 13.

Relationship between average block length and average grain intercept of prior austenite in stir zones of 6%Ni-0.46%C and 6%Ni-0.63%C steels.

The prior austenite grain size in the stir zones was found to decrease significantly as the carbon content increased from 0.30 to 0.46 mass%, and then increased slightly as the carbon content increased from 0.46 mass%. The latter tendency corresponded to the linear relationship between the prior austenite grain size reconstructed from the EBSD data and the carbon content, as shown in Fig. 11. According to Charnack et al.,²⁶⁾ the stacking fault energy (SFE) of plain carbon steels decreases as the carbon content increases up to 0.2 mass%, and then increases as the carbon content increases further. Therefore, it can be inferred that the changes in the prior austenite grain size with the carbon content in the stir zones, as demonstrated in Fig. 12, might be associated with the effect of the carbon content on the SFE.

However, it is known that the packet size and block length are refined as the carbon content increases, although the prior austenite grain sizes are the same. For instance, Kawata et al.²⁷⁾ indicated that in the upper bainite of 9%Ni–x%C steels, the block size decreased as the carbon content increased from 0.15 to 0.30 mass% and did not change significantly as the carbon content further increased from 0.30 mass%. In the martensite structure of 6%Ni–x%C steels investigated in this study, the block length increased slightly as the carbon content further increased from 0.46 mass%, as shown in Fig. 12. Therefore, it is suggested that the change in the block length with the carbon content and the welding condition reflects the change in the prior austenite grain size, particularly in the carbon range exceeding 0.46 mass%. However, it is not appropriate in this study to uniquely evaluate the prior austenite grain size using the block length over the entire range of the carbon contents because the relationship between prior austenite grain size and block length in the carbon range lower than 0.30 mass% is different from that in the range higher than 0.46 mass%. Hence, the average block length was used as a parameter to indicate the prior austenite grain size in the carbon range exceeding 0.46 mass% in the following discussion.

Figure 14 shows the relationship between the average block length and the area fraction of retained austenite in the stir zones of 6%Ni–x_C%C steel. The fraction of retained austenite increased as the prior austenite grain size decreased, regardless of the difference in carbon content and welding condition. This indicates that the increase in the area fraction of retained austenite in the stir zones is significantly correlated with both the carbon content and the refinement of the product martensite block length, namely, the refinement of the prior austenite grain size. This finding is consistent with that of an author’s previous study,¹⁶⁾ in that finer austenite grains exhibited austenite of higher stability, which resulted in a greater amount of retained austenite after FSW. However quantitatively evaluating the effect of the refinement of prior austenite grains on austenite stability is difficult because the effect of the chemical driving force for phase transformation based on the difference in the carbon content is not evaluated separately. Therefore, this aspect is discussed in the following subsection.

Fig. 14.

Relationship between average block length and retained austenite fraction obtained by EBSD in stir zones of 6%Ni-x_C%C steels.

4.2. Effect of Austenite Grain Size on Austenite Stability in Stir Zones

The change in the stability of the retained austenite is likely to be associated with the change in the martensitic transformation behavior. The martensitic transformation start temperature (Ms) is used as the most important parameter to indicate the martensitic transformation behavior. The Ms temperature varies depending on the chemical composition. Herein, the effect of the alloying elements on Ms is expressed as shown in Eq. (1),²⁸⁾ assuming that the content of the alloying element M is x_M mass%.   

$$\begin{array}{l}M\mathrm{s}=521–353x_{\mathrm{C}}–22x_{\mathrm{S}\mathrm{i}}–24.3x_{\mathrm{M}\mathrm{n}}–7.7x_{\mathrm{C}\mathrm{u}}\\ \hspace{1em}\hspace{1em}–17.3x_{\mathrm{N}\mathrm{i}}–17.7x_{\mathrm{C}\mathrm{r}}–25.8x_{\mathrm{M}\mathrm{o}}\end{array}$$(1)

Equation (1) indicates that the Ms temperature decreases by approximately 35°C with a 0.1 mass% increase in the carbon content.

Moreover, the Koistinen–Marburger (KM) model is used to express the relationship between Ms and the amount of retained austenite or martensite. Here, the volume fraction of the retained austenite V_γ at temperature T is expressed as   

$$V_{\gamma }=1-V_{{\alpha }^{\prime }}=exp\left[-{\alpha }_m\left(T_{KM}-T\right)\right],$$(2)

where α_m and T_KM are the transformation rate parameter and the martensite start temperature in the KM model, respectively. Equation (2) can be rewritten as follows:   

$$T_{KM}=T-\frac{ln{V}_{\gamma }}{{\alpha }_m}$$(3)

Once α_m is determined, the martensite start temperature T_KM can be estimated from the volume fraction of retained austenite V_γ after cooling. However, as Eq. (2) does not correspond to the experimental result in the initial stage of transformation until V_α′ reaches 0.15,³⁰⁾ the T_KM calculated using Eq. (3) generally does not equate the Ms temperature measured experimentally. In addition, α_m, whose value is typically between 0.011²⁹⁾ and 0.023,¹⁷⁾ is a parameter associated with the martensite formation rate during cooling. Furthermore, it changes according to the chemical composition and prior austenite grain size. To compensate for the deviation in T_KM from the Ms temperature, the amount of martensite formed during cooling must be measured.¹⁷⁾ Therefore, the correct values of these parameters are difficult to determine.

Hence, in this study, the transformation rate parameter α_m was assumed to be constant, and factors affecting the amount of retained austenite, except for the carbon content, were assumed to emerge as a deviation from the martensitic transformation temperature. Consequently, this apparent deviation in the martensitic transformation temperature was used as an index of austenite stability. This idea implies that α_m was assumed to be the same for all materials, and the apparent KM martensitic transformation start temperature $\overline{T_{KM}}$ was calculated by substituting the area fraction of the retained austenite obtained via EBSD measurement into Eq. (3). The difference between $\overline{T_{KM}}$ and Ms, where Ms is calculated by substituting the carbon content of the base metal into Eq. (1), was defined as $\overline{\mathrm{\Delta }Ms}$ and calculated using Eq. (4).   

$$\overline{\mathrm{\Delta }Ms}=\overline{T_{KM}}-Ms=\frac{ln{V}_{\gamma }}{{\alpha }_m}-T-Ms$$(4)

Here, the correlation between $\overline{\mathrm{\Delta }Ms}$ and the carbon content of the base metal x_C was checked. By adjusting the value of α_m such that the correlation coefficient between $\overline{\mathrm{\Delta }Ms}$ and x_C approached 0, the effect of the carbon content of the base metals was ^judged to be successfully eliminated from $\overline{\mathrm{\Delta }Ms}$. Then, $\overline{T_{KM}}$ and $\overline{\mathrm{\Delta }Ms}$ were calculated by Eq. (4) setting the T to 10°C, which is the room temperature at which FSW was performed, and adjusting the α_m to 0.0132 based on the procedure described above. Ms, $\overline{T_{KM}}$, and $\overline{\mathrm{\Delta }Ms}$ were plotted as functions of x_C and l_b in Fig. 15. As shown in Fig. 15(a), no correlation was observed between $\overline{\mathrm{\Delta }Ms}$ and x_C. However, as shown in Fig. 15(b), $\overline{\mathrm{\Delta }Ms}$ showed a significant correlation with the average block length l_b (i.e., the prior austenite grain size). Hence, the tendency wherein the austenite became more stabilized as the prior austenite grain became finer was clearly demonstrated, as shown in Fig. 14. By combining Figs. 15(b) and 13, it was implied that the decrease in the average intercept length of the prior austenite grains from 2 to 1 μm stabilized the austenite, which corresponded to a reduction of 110°C in the apparent Ms temperature. However, the reduction in the Ms temperature was estimated to be approximately 30°C when extrapolating the decrease in the average intercept length from 2 to 1 μm in the model,³¹⁾ which is validated in the prior austenite grain diameter range greater than 6 μm. Therefore, it is suggested that the calculated $\overline{\mathrm{\Delta }Ms}$ incorporated the effects other than that of the prior austenite grain size, such as the density of the low-angle subgrain boundaries,¹⁵⁾ which increases concurrently with grain refinement. However, because the α_m assumed to be constant in this study is expected to vary depending on the prior austenite grain size, a $\overline{\mathrm{\Delta }Ms}$ value calculated as the difference between Ms and $\overline{T_{KM}}$ temperatures generally does not reflect the behavior of the actual Ms temperature. Therefore, $\overline{\mathrm{\Delta }Ms}$ should be regarded as an index for austenite stability caused by the effects other than that of chemical composition such that an unbiased comparison can be conducted.

Fig. 15.

Relationships between calculated martensitic start temperatures (M_S, $\overline{T_{KM}}$), difference in martensitic start temperatures ($\overline{\mathrm{\Delta }Ms}$) and (a) carbon content of base metal and (b) average block length in stir zones of 6%Ni-x_C%C steels.

As shown in Fig. 15, the value of $\overline{\mathrm{\Delta }Ms}$ is positive for some of the samples and the martensitic transformation temperature seems to be increased. This is because the α_m used in this study was not an actual value, and $\overline{T_{KM}}$ was evaluated to be higher than the actual Ms temperature owing to the underestimation of the retained austenite fractions due to the cut-off of the measurement point with a low CI value.

5. Conclusions

FSW was performed under two welding conditions (i.e., 150 rpm–100 mm/min and 200 rpm–400 mm/min) using 6 mass%Ni steels with different carbon contents ranging from 0.14 to 0.63 mass%. The effects of the carbon content of the base metals and the prior austenite grain size on the amount of retained austenite in the stir zones of the friction stir welded joints were investigated. The maximum temperature attained via FSW was slightly higher in the 150 rpm–100 mm/min, and the cooling rate after FSW was higher in the 200 rpm–400 mm/min.

(1) Irrespective of the welding condition, a microstructure comprising lath martensite and retained austenite was formed in the stir zones of steels with carbon contents exceeding 0.30 mass%.

(2) The amount of retained austenite increased with the carbon content of the base metal.

(3) Greater amounts of retained austenite were formed in the stir zones under the condition of 200 rpm–400 mm/min.

(4) In the stir zones of steels with carbon contents exceeding 0.46 mass%, the average block length of martensite showed a linear relationship with the prior austenite grain size and can be used as an index for the prior austenite grain size.

(5) After excluding the effect of the carbon content, the stability of austenite was evaluated based on the difference in martensitic transformation start temperatures between the apparent $\overline{T_{KM}}$ and Ms. It was revealed that the grain refinement of the prior austenite in the stir zone contributed to the stabilization of austenite.

References

• 1)   W. M.  Thomas,  E. D.  Nicholas,  J. C.  Needhan,  M. G.  Murch,  P.  Temple-Smith and  C. J.  Dawes: Friction Stir Butt Welding, International Patent Application No. PCT/GB92/02203, (1991).
• 2)   Y. S.  Sato,  H.  Kokawa,  K.  Ikeda,  M.  Enomoto,  T.  Hashimoto and  S.  Jogan: Metall. Mater. Trans. A, 32 (2001), 941. https://doi.org/10.1007/s11661-001-0351-z
• 3)   M. W.  Mahoney,  C. G.  Rhodes,  J. G.  Flintoff,  W. H.  Bingel and  R. A.  Spurling: Metall. Mater. Trans. A, 29 (1998), 1955. https://doi.org/10.1007/s11661-998-0021-5
• 4)   H. J.  Liu,  H.  Fujii,  M.  Maeda and  K.  Nogi: J. Mater. Process. Technol., 142 (2003), 692. https://doi.org/10.1016/S0924-0136(03)00806-9
• 5)   H.  Fujii,  L.  Cui,  M.  Maeda and  K.  Nogi: Mater. Sci. Eng. A, 419 (2006), 25. https://doi.org/10.1016/j.msea.2005.11.045
• 6)   A. P.  Reynolds,  W.  Tang,  T.  Gnaupel-Herold and  H.  Prask: Scr. Mater., 48 (2003), 1289. https://doi.org/10.1016/S1359-6462(03)00024-1
• 7)   Y. S.  Sato,  T. W.  Nelson and  C. J.  Sterling: Acta Mater., 53 (2005), 637. https://doi.org/10.1016/j.actamat.2004.10.017
• 8)   L.  Cui,  H.  Fujii,  N.  Tsuji,  K.  Nakata,  K.  Nogi,  R.  Ikeda and  M.  Matsushita: ISIJ Int., 47 (2007), 299. https://doi.org/10.2355/isijinternational.47.299
• 9)   Z.  Wu,  T.  Nagira,  K.  Ushioda,  G.  Miyamoto and  H.  Fujii: Mater. Sci. Eng. A, 799 (2021), 140068. https://doi.org/10.1016/j.msea.2020.140068
• 10)   T.  Kawakubo,  T.  Nagira,  K.  Ushioda and  H.  Fujii: ISIJ Int., 61 (2021), 2150. https://doi.org/10.2355/isijinternational.ISIJINT-2021-007
• 11)   R.  Rai,  A.  De,  H. K. D. H.  Bhadeshia and T. DebRoy: Sci. Technol. Weld. Join., 16 (2011), 325. https://doi.org/10.1179/1362171811Y.0000000023
• 12)   T.  Miura,  R.  Ueji,  H.  Fujii,  H.  Komine and  J.  Yanagimoto: Mater. Des., 90 (2016), 915. https://doi.org/10.1016/j.matdes.2015.11.037
• 13)   T.  Miura,  R.  Ueji,  H.  Fujii,  T.  Murakami and  T.  Kobashi: Sci. Technol. Weld. Join., 21 (2016), 281. https://doi.org/10.1080/13621718.2015.1104096
• 14)   H.  Fujii,  R.  Ueji,  Y.  Morisada and  H.  Tanigawa: Scr. Mater., 70 (2014), 39. https://doi.org/10.1016/j.scriptamat.2013.09.012
• 15)   T.  Miura,  R.  Ueji and  H.  Fujii: J. Mater. Process. Technol., 216 (2015), 216. https://doi.org/10.1016/j.jmatprotec.2014.09.014
• 16)   T.  Miura,  R.  Ueji and  H.  Fujii: Weld. Int., 32 (2018), 338. https://doi.org/10.1080/09507116.2017.1346831
• 17)   C.  Celada-Casero,  J.  Sietsma and  M. J.  Santofimia: Mater. Des., 167 (2019), 107625. https://doi.org/10.1016/j.matdes.2019.107625
• 18)   A. D.  Romig and  J. I.  Goldstein: Metall. Trans. A, 9 (1978), 1599. https://doi.org/10.1007/BF02661942
• 19)   V.  Manvatkar,  A.  De,  L.-E.  Svensson and T. DebRoy: Scr. Mater., 94 (2015), 36. https://doi.org/10.1016/j.scriptamat.2014.09.012
• 20)   L.  Cheng,  A.  Böttger,  Th. H.  de Keijser and  E. J.  Mittemeijer: Scr. Metall. Mater., 24 (1990), 509. https://doi.org/10.1016/0956-716X(90)90192-J
• 21)   G.  Miyamoto,  N.  Iwata,  N.  Takayama and  T.  Furuhara: Acta Mater., 58 (2010), 6393. https://doi.org/10.1016/j.actamat.2010.08.001
• 22)   M.  Abbasi,  T. W.  Nelson,  C. D.  Sorensen and  L.  Wei: Mater. Charact., 66 (2012), 1. https://doi.org/10.1016/j.matchar.2012.01.010
• 23)   S.  Morito,  H.  Tanaka,  R.  Konishi,  T.  Furuhara and  T.  Maki: Acta Mater., 51 (2003), 1789. https://doi.org/10.1016/S1359-6454(02)00577-3
• 24)   S.  Morito,  X.  Huang,  T.  Furuhara,  T.  Maki and  N.  Hansen: Acta Mater., 54 (2006), 5323. https://doi.org/10.1016/j.actamat.2006.07.009
• 25)   T.  Maki,  K.  Tsuzaki and  I.  Tamura: Tetsu-to-Hagané, 65 (1979), 515 (in Japanese). https://doi.org/10.2355/tetsutohagane1955.65.5_515
• 26)   W.  Charnock and  J.  Nutting: Met. Sci. J., 1 (1967), 123. https://doi.org/10.1179/msc.1967.1.1.123
• 27)   H.  Kawata,  K.  Sakamoto,  T.  Moritani,  S.  Morito,  T.  Furuhara and  T.  Maki: Mater. Sci. Eng. A, 438–440 (2006), 140. https://doi.org/10.1016/j.msea.2006.02.175
• 28)  Welding Technology Information Center of The Japan Welding Engineering Society: Calculation Software of Welding, The Japan Welding Engineering Society - The Welding Information Center, http://www-it.jwes.or.jp/weld_simulator/cal1.jsp, (accessed 2021-09-30), (in Japanese).
• 29)   D. P.  Koistinen and  R. E.  Marburger: Acta Metall., 7 (1959), 59. https://doi.org/10.1016/0001-6160(59)90170-1
• 30)   S. M. C.  van Bohemen and  J.  Sietsma: Mater. Sci. Technol., 30 (2014), 1024. https://doi.org/10.1179/1743284714Y.0000000532
• 31)   A.  Arlazarov,  E.  Soares Barreto,  N.  Kabou and  D.  Huin: Metall. Mater. Trans. A, 51 (2020), 6159. https://doi.org/10.1007/s11661-020-06022-6