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Microstructure feature of friction stir butt-welded ferritic ductile iron Hung-Tu Chang a , Chaur-Jeng Wang a , Chin-Pao Cheng b,a Department of Mechanical Engineering, National Taiwan University of Science and Technology, Taipei, Taiwan, ROC b Department of Mechatronic Technology, National Taiwan Normal University, Taipei, Taiwan, ROC article info Article history: Received 27 August 2013 Accepted 19 November 2013 Available online 4 December 2013 Keywords: Friction stir welding Ductile iron Dynamic recrystallization Martensite abstract This study conducted friction stir welding (FSW) by using the butt welding process to join ferritic ductile iron plates and investigated the variations of microsturcture in the joined region formed after welding. No defects appeared in the resulting experimental weld, which was formed using a 3-mm thick ductile iron plate and tungsten carbide alloy stir rod to conduct FSW at a rotational speed of 982 rpm and trav- eling speed of 72 mm/min. The welding region was composed of deformed graphite, martensite phase, and dynamically recrystallized ferrite structures. In the surface region and on the advancing side (AS), the graphite displayed a striped configuration and the ferritic matrix transformed into martensite. On the retreating side (RS), the graphite surrounded by martensite remained as individual granules and the matrix primarily comprised dynamically recrystallized ferrite. After welding, diffusion increased the carbon content of the austenite around the deformed graphite nodules, which transformed into mar- tensite during the subsequent cooling process. A micro Vickers hardness test showed that the maximum hardness value of the martensite structures in the weld was approximately 800 HV. An analysis using an electron probe X-ray microanalyzer (EPMA) indicated that its carbon content was approximately 0.7–1.4%. The peak temperature on the RS, 8 mm from the center of the weld, measured 630 °C by the thermocouple. Overall, increased severity of plastic deformation and process temperature near the upper stir zone (SZ) resulted in distinct phase transformation. Furthermore, the degree of plastic deformation on the AS was significantly greater than that on the RS, and relatively complete graphite granules and the fine ferrite grains resulting from dynamic recrystallization were observed on the RS. Ó 2013 Elsevier Ltd. All rights reserved. 1. Introduction Friction stir welding (FSW) is a novel solid-state welding technology that adopts the heat generated by the friction that re- sults from a stir rod rotating at high speed and penetrating a base material, subsequently achieving the joining effect during the feeding process through the metal plastic flow phenomenon cre- ated by stirring [1–8]. During the welding process, the temperature remains below the melting point of the welding material. In contrast to traditional welding methods, the absence of melting renders FSW applicable to materials that are difficult to weld such as age-hardened aluminum alloys [9–11]. Several recent studies have employed FSW to join carbon steel and stainless steel [4–8,12–20]. The rotating tools must tolerate high temperatures (1000 °C), which are required to create plastic flow when welding iron-based alloys. The complex phase transformation of steel materials varies substantially with the carbon content of the base material. Furthermore, temperature changes that occur during welding and the cooling rate after welding significantly influence the microstructure of the welding regions, altering the mechanical properties of weldments [4,5,16–20]. Fujii et al. performed FSW on 0.12% carbon steel by using welding parameters of 400 rpm and 100 mm/min, and observed relatively small ferrite grains (3 lm) that resulted from dynamic recrystallization at the bottom of the weld [16]. Furthermore, Fujii et al. extensively investigated how welding temperature affects structural changes in various weld regions. By using 0.85% carbon tool steel for FSW, Chung et al. reported that martensite structures formed in the weld when the welding temperature exceeded that at the eutectoid transforma- tion point (A 1 ) [17]. A welding temperature lower than that of the A 1 transformation temperature prevented martensite forma- tion in the ferrite matrix and simultaneously induced grain refine- ment; the weldment also demonstrated superior toughness and ductility. Because Chung et al. exercised joint material which is a ferrite matrix with globular cementite, the structure characteristic obtained proper plastic flow in lower temperature by FSW. Conse- quently, the base material was joined by friction stir effect, and without any phase transformation produced in the joint region. Sun et al. indicated that performing FSW using a 3.2-mm thick, 0.45% carbon steel at a rotational speed of 600 rpm and a traveling speed greater than 300 mm/min produces martensite and bainite structures in the weld; however, a traveling speed lower than 0261-3069/$ - see front matter Ó 2013 Elsevier Ltd. All rights reserved. http://dx.doi.org/10.1016/j.matdes.2013.11.046 Corresponding author. Address: 162, HePing East Road, Section 1, Taipei, Taiwan, ROC. Tel.: +886 2 77343521; fax: +886 2 23583074. E-mail address: [email protected] (C.-P. Cheng). Materials and Design 56 (2014) 572–578 Contents lists available at ScienceDirect Materials and Design journal homepage: www.elsevier.com/locate/matdes

Microstructure feature of friction stir butt-welded ferritic ductile iron

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Materials and Design 56 (2014) 572–578

Contents lists available at ScienceDirect

Materials and Design

journal homepage: www.elsevier .com/locate /matdes

Microstructure feature of friction stir butt-welded ferritic ductile iron

0261-3069/$ - see front matter � 2013 Elsevier Ltd. All rights reserved.http://dx.doi.org/10.1016/j.matdes.2013.11.046

⇑ Corresponding author. Address: 162, HePing East Road, Section 1, Taipei,Taiwan, ROC. Tel.: +886 2 77343521; fax: +886 2 23583074.

E-mail address: [email protected] (C.-P. Cheng).

Hung-Tu Chang a, Chaur-Jeng Wang a, Chin-Pao Cheng b,⇑a Department of Mechanical Engineering, National Taiwan University of Science and Technology, Taipei, Taiwan, ROCb Department of Mechatronic Technology, National Taiwan Normal University, Taipei, Taiwan, ROC

a r t i c l e i n f o

Article history:Received 27 August 2013Accepted 19 November 2013Available online 4 December 2013

Keywords:Friction stir weldingDuctile ironDynamic recrystallizationMartensite

a b s t r a c t

This study conducted friction stir welding (FSW) by using the butt welding process to join ferritic ductileiron plates and investigated the variations of microsturcture in the joined region formed after welding.No defects appeared in the resulting experimental weld, which was formed using a 3-mm thick ductileiron plate and tungsten carbide alloy stir rod to conduct FSW at a rotational speed of 982 rpm and trav-eling speed of 72 mm/min. The welding region was composed of deformed graphite, martensite phase,and dynamically recrystallized ferrite structures. In the surface region and on the advancing side (AS),the graphite displayed a striped configuration and the ferritic matrix transformed into martensite. Onthe retreating side (RS), the graphite surrounded by martensite remained as individual granules andthe matrix primarily comprised dynamically recrystallized ferrite. After welding, diffusion increasedthe carbon content of the austenite around the deformed graphite nodules, which transformed into mar-tensite during the subsequent cooling process. A micro Vickers hardness test showed that the maximumhardness value of the martensite structures in the weld was approximately 800 HV. An analysis using anelectron probe X-ray microanalyzer (EPMA) indicated that its carbon content was approximately0.7–1.4%. The peak temperature on the RS, 8 mm from the center of the weld, measured 630 �C by thethermocouple. Overall, increased severity of plastic deformation and process temperature near the upperstir zone (SZ) resulted in distinct phase transformation. Furthermore, the degree of plastic deformation onthe AS was significantly greater than that on the RS, and relatively complete graphite granules and thefine ferrite grains resulting from dynamic recrystallization were observed on the RS.

� 2013 Elsevier Ltd. All rights reserved.

1. Introduction

Friction stir welding (FSW) is a novel solid-state weldingtechnology that adopts the heat generated by the friction that re-sults from a stir rod rotating at high speed and penetrating a basematerial, subsequently achieving the joining effect during thefeeding process through the metal plastic flow phenomenon cre-ated by stirring [1–8]. During the welding process, the temperatureremains below the melting point of the welding material. Incontrast to traditional welding methods, the absence of meltingrenders FSW applicable to materials that are difficult to weld suchas age-hardened aluminum alloys [9–11]. Several recent studieshave employed FSW to join carbon steel and stainless steel[4–8,12–20]. The rotating tools must tolerate high temperatures(1000 �C), which are required to create plastic flow when weldingiron-based alloys. The complex phase transformation of steelmaterials varies substantially with the carbon content of the basematerial. Furthermore, temperature changes that occur duringwelding and the cooling rate after welding significantly influence

the microstructure of the welding regions, altering the mechanicalproperties of weldments [4,5,16–20]. Fujii et al. performed FSW on0.12% carbon steel by using welding parameters of 400 rpm and100 mm/min, and observed relatively small ferrite grains (3 lm)that resulted from dynamic recrystallization at the bottom of theweld [16]. Furthermore, Fujii et al. extensively investigated howwelding temperature affects structural changes in various weldregions. By using 0.85% carbon tool steel for FSW, Chung et al.reported that martensite structures formed in the weld when thewelding temperature exceeded that at the eutectoid transforma-tion point (A1) [17]. A welding temperature lower than that ofthe A1 transformation temperature prevented martensite forma-tion in the ferrite matrix and simultaneously induced grain refine-ment; the weldment also demonstrated superior toughness andductility. Because Chung et al. exercised joint material which is aferrite matrix with globular cementite, the structure characteristicobtained proper plastic flow in lower temperature by FSW. Conse-quently, the base material was joined by friction stir effect, andwithout any phase transformation produced in the joint region.Sun et al. indicated that performing FSW using a 3.2-mm thick,0.45% carbon steel at a rotational speed of 600 rpm and a travelingspeed greater than 300 mm/min produces martensite and bainitestructures in the weld; however, a traveling speed lower than

Fig. 1. Schematic illustration of friction stir butt-welding.

Table 1Original size of welding tool, mm.

Tool material Shoulder diameter Pin diameter Pin length

WC-Mo 12 3.6 2.8

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300 mm/min inhibits the formation of martensite and bainite [5].Lakshminarayanan et al. identified very fine ferrite and martensitestructures in the weld when performing FSW at 1000 rpm and50 mm/min and using ferrite-based stainless steel, which contains0.026% carbon, 11.4% chromium, and 0.4% nickel [6]. According tothe aforementioned studies, the chemical composition of iron-based alloys and changes in welding temperature extensively alterthe microstructure of welds, affecting the mechanical properties ofweldments.

Compared to traditional fusion welding technologies, FSWprovides a feasible method for joining ductile irons. Previously,welding methods were rarely adopted for joining ductile irons be-cause of various problems that could not be overcome during thewelding process. The carbon content of ductile iron is considerablyhigher than that of carbon steel, causing poor weldability; thus,during high-temperature fusion welding, carbon in the graphitedissolves into the surrounding melted region and diffuses acrossthe unmelted regions that contain austenite phase. Consequently,the hard and brittle carbide and martensite developed during thecooling and solidifying processes reduce the mechanical propertiesof the weldments [2,21,22]. To resolve this problem, several stud-ies have employed FSW to join ductile irons. By conducting dissim-ilar welding using ductile iron and low-carbon steel, our previousstudies observed martensite and pearlite structures in the weld,and found that proper heat treatment removes martensite andimproves the tensile properties of the weldments [2,3]. Fujiiet al. and Imagawa et al. conducted friction stir processing withferritic ductile iron, pearlitic ductile iron and flake graphite castiron for surface hardening processing that rendered the hardeningeffect by forming martensite transformation in the stir zone. Theirstudies indicated that the carbon atom is more difficult to diffusein the ferritic matrix of ductile iron, so that the optimal conditionof friction stir processing was narrower than that of other twomaterials [23–25]. Furthermore, Cheng et al. applied friction stirsurface hardening to ductile iron by using a stir rod without apin and rotating and traveling speeds of 2200 rpm and 60 mm/min, respectively. Mixtures of structures containing ferrite, bainite,martensite, and retained austenite were found in the stirred region,and this hardened layer improved the resistance of ductile iron toerosion [26]. Based on previous studies, martensite forms in thejoined region of ductile iron after FSW; however, the mechanismrelated to this formation process has not been investigated. There-fore, this study employs FSW to join ductile iron by using the buttwelding technique, investigating the morphological and micro-structural changes in the graphite surrounding the weld.

2. Experimental procedure

The base material in the current experiment comprised a 3-mmthick ferritic ductile iron plates that contained 2.0% carbon, 2.5%silicon, 0.09% manganese, 0.006% sulfur, 0.034% phosphorous,0.039% magnesium, and iron. Based on the experimental require-ments, the ductile iron plate was processed to the desired dimen-sions (95 mm � 40 mm � 3 mm). Before welding, the oxidizedlayer on the ductile iron surface was sanded, washed with acetone,and blow-dried. Fig. 1 presents a schematic of the relative positionsof the base material and stir rod during the butt welding process.Furthermore, Table 1 lists the dimensions of the stir rod, whichequips with columnar probe without threads. During the weldingprocess, the stir rod was tilted 0.5 degrees from the normal direc-tion of the plate. To determine the optimal parameters for thewelding process, pilot tests were performed by different rotationaland travelling speed. The preliminary results showed thatrotational and traveling speeds of 982 rpm and 72 mm/min,respectively, were favorable for joining. After welding, the weld

surface was smooth and lacked defects, the ductile iron platewas not deformed, and the rotating tool (i.e., stir rod) presentedno wear. To explore how temperature changes near the weld af-fected the microstructure of the ductile iron, the K-type thermo-couples were installed near the retreating side (RS) 8, 13 and18 mm from the center of the weld. These thermocouples wereembedded into three holes, each 5 mm apart, to record the temper-ature changes that occurred during welding. Subsequently, a sec-tion of the weld region was sliced to facilitate an observation ofits microstructure, and the specimen was polished before beingetched using Nital etchant. The microstructural changes in theweld of the etched specimen were observed using an opticalmicroscope and scanning electron microscope (SEM).

After FSW, variation in the hardness of the joined region wasfirst measured by polishing the cross-section of the sliced speci-men and then conducting micro Vickers hardness tests, whichwere performed as per ASTM: E384-11e1 specification. Themeasurement conditions were a load of 200 g and a hold time of15 s. To determine the variation in hardness values of differentregions, a microhardness distribution curve was plotted using thedata obtained from the following procedure: dividing the cross-section of the specimen into three regions (surface, middle, andbottom) from top to bottom and using the welding line as the cen-ter, microhardness was measured every 0.2 mm on each side of thecenter point. Additionally, to verify the phenomenon by which car-bon atoms diffuse from graphite during welding, the current studyadopted JEOL JXA-8200 electron probe X-ray microanalyzer(EPMA) for quantitative and line scan analyses. In addition to theEMPA results, the microhardness values and phase transformationresults were employed to explore potential factors that caused theformation of various weld phases and structures.

3. Results and discussion

3.1. Microstructural characterization

Fig. 2 shows the microstructure of ductile iron used as the FSWbase material. The graphite nodules (approximately 20–25 lm indiameter) were evenly distributed within the ferrite matrix, whichwere in the state of equiaxed crystals approximately 30–40 lm indiameter. Fig. 3 presents a cross-sectional schematic of the weldregions. The sub-surface region of the weld is the area where thestir rod shoulder acts on the base material and is approximatelythe width of the stir rod shoulder. The spindle rotation createsfriction between the shoulder and base material; thus, the location

Fig. 2. Microstructures of the ductile iron base metal.

Fig. 3. Schematic diagram of the welded nugget and part names. (A: surface region;B: middle SZ; C: bottom SZ; D: advanced side TMAZ; E: retreated side TMAZ).

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where the center of the surface region overlaps the stir zone (SZ)has the highest temperature [7]. The center and bottom of the SZare regions in which the stir rod pin directly acts on the material.The thermo-mechanically affected zone (TMAZ) is located outsidethe SZ and is divided into the advancing side (AS) and the retreat-ing side (RS) based on the difference of the rotating direction of stirrod and the feeding direction of base material. Because the basematerial is composed of as-cast material, the heat-affected zone(HAZ) in the current study was indistinct. The area outside theTMAZ is the base material region that is unaffected by the weldingprocess. In Fig. 3, region D represents the AS of the TMAZ (AS-T-MAZ), and E represents the RS of the TMAZ (RS-TMAZ). The micro-structures of various locations (from the surface to the bottomregions) in the weld cross-section demonstrate significant differ-ences because of varying temperature effects.

The cross-section of the weld region in Fig. 4 shows a successjoining effect, exhibiting no holes or other defects. In addition, sig-nificant structural differences exist between the AS and RS of theweld. The graphite granules on the RS remained granular-shaped,but those on the AS and in the surface region demonstrated adistinct black-striped configuration. This phenomenon occurredbecause during the FSW process, the plastic flow resulting fromstress stretched the graphite nodules, resulting in striped patternsthat reflect the tracks of the plastic flow [2]. According to previous

Fig. 4. The macrostructure of the welded nugget.

studies, the graphite nodules in ductile iron can be consideredpores that are unaffected by stress. When plastic deformation oc-curs in stressed ferrite matrix, graphite acts as a packing materialthat fills the pores in the ferrite matrix [27]; thus, graphite exhibitsthe plastic flow path of ferritic matrix.

Fig. 5 presents a magnified view of the various welding regionspresented in Fig. 4. Fig. 5(a)–(c) shows that because of the plasticflow generated by the rotating stir rod shoulder, the graphitenodules in the surface region were scattered into a striped config-uration, and numerous dense martensite structures were formed.In Fig. 5(a) and (b), the martensite in the surface region is rela-tively dense, displaying an unparalleled arrangement and distinctmidrib characteristics. According to Krauss, plate martensite isformed when high-carbon or alloy steels are quenched [28];therefore, the martensite structures in the surface region possesshigh carbon content. Imagawa et al. conducted friction stir pro-cessing with ferritic ductile iron at rotational speed of 900 rpmand traveling speed 50, 100 mm/min. In the stir zone, the sphe-roidal graphite was crushed and striated. In addition, martensitestructure was formed in the matrix [24]. Fig. 5(c) indicates the RSof the surface region. Approximately 5- to 10-lm fine ferritegrains are present within the graphite, and the distribution ofmartensite in this region is noticeably low. Chung et al. indicatedthat martensite forms in regions in which the processing temper-ature is higher than that at the transformation point A1 whenFSW is conducted using 0.85% carbon steel [17]. Based on themartensite structures that formed in the surface region, the tem-perature in this region exceeded the temperature at the transfor-mation point A1 during FSW. Fig. 5(d)–(f) shows the mixture ofmartensite and ferrite structures identified in the middle region.In Fig. 5(d) and (e), the proportion of martensite in the AS inter-face and middle SZ region remains high. Fig. 5(f) presents themiddle region of the RS-TMAZ, in which graphite nodules in themicrostructure remained granular-shaped and are surroundedby martensite and fine ferrite grains. Furthermore, ferrite struc-tures were primarily observed in the microstructure of the bot-tom region. The AS-TMAZ (Fig. 5(g)) and bottom SZ region(Fig. 5(h)) demonstrated partial phase transformation. The RSmatrix (Fig. 5(i)) was primarily composed of ferrite grains formedby dynamic recrystallization. In summary, as the severe plasticdeformation (SPF) near the upper SZ region increases, the plasticdeformation on the AS is greater than that on the RS, therebyexhibiting distinct phenomena of phase transformation andgraphite deformation.

3.2. Hardness distribution

Fig. 6 shows the Vickers hardness profile of the ductile ironcross-section following FSW. The hardness of the SZ, TMAZ, andbase material region was measured. The hardness value of the duc-tile iron base material was approximately 160–180 HV. Accordingto the microhardness curve of the surface region, the maximumhardness value reached 800 HV because numerous martensitestructures were observed in this region. Regarding the surface re-gion, the hardness value of the SZ was the highest, followed by thatof the AS-TMAZ. The hardness value of the RS-TMAZ was slightlylower than that of the AS, and the hardness of the base materialwas the lowest. These results show that severe stirring effect anda temperature greater than that at the transformation point A1 en-hances the formation of martensite phase. Krauss determined thatthe hardness value of the martensite phase was corresponding tocarbon content [28]. Consequently, the hardness value for themicrostructure of the surface SZ region (Fig. 5(b)) exceeded800 HV, suggesting that the carbon content of the martensitestructures was relatively high. A small quantity of martensitewas observed in the middle region of the SZ (Fig. 5(e)), and the

Fig. 5. The microstructure of the welded nugget: (a)–(i) microstructure of the marked regions.

Fig. 6. Microhardness curve of the ductile iron welded at a rotational speed of982 rpm and traveling speed of 72 mm/min.

Fig. 7. Thermal cycle curve of the ductile iron welded at a rotational speed of982 rpm and traveling speed of 72 mm/min.

H.-T. Chang et al. / Materials and Design 56 (2014) 572–578 575

hardness value of this region was approximately 300–350 HV.According to Chung, various peak temperatures influence the dis-tribution and quantity of martensite, rendering the hardness valueof the middle SZ region lower than that of the surface region [17].Regarding the bottom SZ region, because the peak temperatureduring welding was low, a eutectoid transformation temperaturewas not achieved and martensite structures were not formed.The hardness value of this region ranged between 250 and300 HV, which remained higher than that of the base material.Numerous studies have indicated that when carbon steel is usedin FSW, refined grains are generated in the SZ because of thedynamic recrystallization caused by severe plastic deformation athigh temperatures [2,6,16–20]. The structures of these refinedgrains comprised subgrain boundaries and dislocations; thus,refined grains increase hardness values [16].

To examine the temperature changes near the weld during FSW,the K-type thermocouples were adopted to measure and recordthermal cycles during welding (Fig. 7). The welding parametersused in the experiment included a rotational speed of 982 rpmand a traveling speed of 72 mm/min. The peak temperature in-creased when measured nearer the SZ, yielding a peak temperatureof 630 �C on the RS, 8 mm from the center of the SZ. The peaktemperature of the surface region near the center of the weldwas estimated to be higher than the eutectoid transformationtemperature. Sun et al. conducted FSW using 0.45% carbon steel,

Fig. 8. Relatively small grains in the welded nugget: (a) microstructure of thebottom SZ region; (b) microstructure of the RS-TMAZ; and (c) microstructure of theRS-TMAZ, captured using FE-SEM.

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and reported that the peak temperature during welding measuredby the thermocouple embedded in the stir rod shoulder was higherthan 900–1000 �C [5]. Zhu et al. performed welding on 304L stain-less steel using parameters of 500 rpm and 100 mm/min, anddetermined that the temperature of the region 14 mm from theweld center was approximately 480 �C; according to a simulation,they estimated that the temperature of the region where the weldand stir rod shoulder interacted was between 900 �C and 1157 �Cor higher [8]. In addition, the shoulder of the stir rod used duringFSW demonstrated geometric heat-collecting capabilities, render-ing the temperature of the upper part of the weld higher than thatof the lower part, and the AS temperature greater than the RS tem-perature [7]. Because the joint mechanism of FSW was material’splastic flow in high temperature ambient, which make the weldsachieve the joining effect. Therefore comparing and analyzing thetemperatures obtained in the current experiment with those ofother studies, requiring the weld to have a peak temperature of900–1000 �C or higher is reasonable for achieving superior welda-bility of iron-based alloys. While some studies indicated that thehigh carbon steel can be welded soundly below eutectoid temper-ature by FSW, and avoided producing martensite structure in thewelding nugget, this characteristic was attributed to a greatamount of continuous plastic flow in the welding region whichcould be created in lower temperature [16–19]. In this study, thematrix of ductile iron was ferritic structure, but the graphite gran-ules distributing in the matrix without ductility degraded thedeformation plastic flow in high temperature. Consequently, themetal plates required higher temperature to obtain sound welding.

Fig. 8 shows that the fine grains in the bottom SZ region of theweld and the middle region of the RS-TMAZ demonstrate two mor-phologies. Fig. 8(a) presents the microstructure of the bottom SZregion, which is the area where the stir rod pin and ductile ironbase material interact. Because the peak temperature of the bot-tom SZ region was lower than that of the surface or middle regions,the grain size was approximately 4–5 lm and the grains in this re-gion exhibited incomplete growth. Fig. 8(b) shows the fine grainstructures in the RS-TMAZ, and Fig. 8(c) indicates the microstruc-ture of the same region observed using FE-SEM. In contrast tothe microstructure in Fig. 8(a), the fine grains in Fig. 8(c) is sur-rounded by martensite structures and approximately 1–4 lm indiameter. Based on the phase transformation temperature for mar-tensite, the temperature of the RS-TMAZ during welding shouldhave reached the eutectoid temperature. However, because theperiod during which temperature is maintained above the trans-formation temperature was extremely short, only the structuresaround the graphite transformed into martensite during cooling.Moreover, the fine ferrite grains that resulted from the SPD duringrotation were influenced by the surrounding martensite structures,which restricted recrystallization and grain growth. In other words,the ferrite in the RS-TMAZ was not subjected to a transformationtemperature, and the carbon content failed to reach the levelrequired for austenite formation.

3.3. EPMA analysis

To examine the diffusion of carbon atoms in the ductile ironduring FSW, four regions were selected for a quantitative analysiswith an electron probe micro-analyzer (EPMA) (Fig. 9(a)). LocationA represents the region that contains martensite and a carbon con-tent of 0.727%; B represents the region around the graphite thatcontains martensite and a carbon content of 1.02%; C and D areindicated on the lower-right corner of the diagram and enlargedin Fig. 9(b), where C is the grain boundaries comprising a carboncontent of 1.421% and D is the grain interior that possesses acarbon content of 0.116%. Table 2 summarizes the compositionalpercentages of iron, carbon, silicon, and manganese present in

the four observed regions. Fig. 10 shows the carbon content curveobtained using EPMA line scan analysis; the white regions in thefigure are the grain boundaries where martensite may form. Basedon the EPMA analysis, the relationship between the carbon contentof the martensite within the weld and microhardness corre-sponded to that obtained by Krauss [28]. In other words, martens-ite with a hardness value exceeding 800 HV exhibits carboncontent greater than 1.0%. Moreover, the EPMA line scan data showthat the carbon content of ferrite grains was lower than that of thegrain boundaries.

At normal temperature, the solubility of carbon atoms in fer-rites is extremely low. When temperatures exceed the eutectoidtransformation temperature, the ferrite structures around graphitetransform into austenite, which has higher carbon solubility. Insuch conditions, the carbon atoms in graphite nodules graduallydiffuse, forming carbon-rich austenite [29–31]. However, whenductile iron was subjected to FSW, the weld temperature wasmaintained above the transformation temperature for a short per-iod, impeding the formation of evenly distributed and carbon-richaustenite. Consequently, austenite formed only around the graph-ite because the distance of diffusion was relatively short and thecarbon content was comparatively high in this region. In other

Fig. 9. EPMA diagram: (a) regions in which quantitative analysis was performedusing EPMA; and (b) points C and D.

Table 2Composition derived from EPMA quantitative analysis.

Region Fe% C% Si% Mn% Total%

A 94.744 0.727 4.399 0.130 100B 94.090 1.020 4.778 0.112 100C 94.247 1.421 4.188 0.144 100D 95.548 0.116 4.277 0.059 100

Fig. 10. EPMA line scan curve of carbon elements.

Fig. 11. Relationship between the carbon content and distance from graphitenodules (a) carbon content gradient of ductile iron before FSW; and (b) carboncontent gradient of ductile iron after FSW.

H.-T. Chang et al. / Materials and Design 56 (2014) 572–578 577

words, the carbon content in the matrix and the distance fromgraphite nodules exhibited an inverse relationship. Fig. 11 presentsthe curves of the carbon content gradient in differing regions of theductile iron following FSW. Because temperature influences thediffusion behavior of carbon atoms, the high operative temperatureof the surface region increased the carbon content in the matrix,the distance by which carbon atoms diffused, and the quantity ofmartensite formed during cooling. Conversely, the operative tem-peratures in the middle region and on the RS were relativelylow; therefore, martensite formed only around the graphite nod-ules during cooling.

4. Conclusions

The following findings were obtained based on the experimen-tal results:

(1) Welding ductile iron by employing FSW at a rotational speedof 982 rpm and traveling speed of 72 mm/min produces aweld that displays a smooth surface and no defects.

(2) The weld microstructure is composed of deformed graphite,martensite, and dynamically recrystallized ferrite. High tem-peratures and the friction resulting from rotation caused thecarbon atoms to diffuse from the graphite into matrix, trans-forming certain regions into martensite and other regionsinto dynamic recrystallized ferrite. The microhardness testand EPMA analysis showed that the carbon content of themartensite was approximately 0.7–1.4%, and the maximumhardness value was 800 HV.

(3) The graphite in the surface region and on the AS exhibited astriped configuration, and distinct martensite structuresformed in the matrix. The graphite on the RS remained asindividual granules, and martensite was observed outsidethe graphite. The matrix was primarily composed of dynam-ically recrystallized ferrite. The proportion of martensite inthe weld surface region and AS matrix was significantlyhigher than those on the RS and in the bottom region ofthe weld.

References

[1] Thomas WM, Nicholas ED, Needham JC, Murch MG, Temple-smith P, Dawes CJ.international patent application no. PCT/GB92/02203.

[2] Cheng CP, Lin HM, Lin JC. Friction stir welding of ductile iron and low carbonsteel. Sci Technol Weld Join 2010;15:706–11.

[3] Chang HT, Wang CJ, Cheng CP. Friction stir lap welded low carbon steel andductile iron: microstructure and mechanical properties. Sci Technol Weld Join2013;18:688–96.

578 H.-T. Chang et al. / Materials and Design 56 (2014) 572–578

[4] Cho HH, Kang SH, Kim SH, Oh KH, Kim HJ, Chang WS, et al. Microstructuralevolution in friction stir welding of high-strength linepipe steel. Mater Des2012;34:258–67.

[5] Sun YF, Konishi Y, Kamai M, Fujii H. Microstructure and mechanical propertiesof S45C steel prepare by laser-assisted friction stir welding. Mater Des2013;47:842–9.

[6] Lakshminarayanan AK, Balasubramanian V. An assessment of microstructure,hardness, tensile and impact strength of friction stir welded ferritic stainlesssteel joints. Mater Des 2010;31:4592–600.

[7] Cho JH, Boyce DE, Dawson PR. Modeling strain hardening and textureevolution in friction stir welding of stainless steel. Mater Sci Eng A2005;398:146–63.

[8] Zhu XK, Chao YJ. Numerical simulation of transient temperature and residualstresses in friction stir welding of 304L stainless steel. J Mater Process Technol2004;146:263–72.

[9] Sato YS, Kokawa H, Enmoto M, Jogan S, Hashimoto T. Microstructural evolutionof 6063 aluminum during friction-stir welding. Metall Mater Trans A1999;30:2429–37.

[10] Rhodes CG, Mahoney MW, Bingel WH, Spurling RA, Bampton CC. Effects offriction stir welding on microstructure of 7075 aluminum. Scripta Mater1997;36:69–75.

[11] Su JQ, Nelson TW, Mishra R, Mahoney M. Microstructure investigation offriction stir welded 7050–T651 aluminum. Acta Mater 2003;51:713–29.

[12] Fazel-Najafabadi M, Kashani-Bozorg SF, Zarei-Hanzaki A. Joining of CP-Ti to304 stainless steel using friction stir welding technique. Mater Des2010;31:4800–7.

[13] Thomas WM, Threadgill PL, Nicholas ED. Feasibility of friction stir weldingsteel. Sci Technol Weld Join 1999;4:365–72.

[14] Reynolds AP, Tang W, Posada M, DeLoach J. Friction stir welding of DH 36 steel.Sci Technol Weld Join 2003;8:455–61.

[15] Lienert TJ, Stellwag Jr WL, Grimmett BB, Warke RW. Friction stir weldingstudies on mild steel. Weld J 2003:1–9.

[16] Fujii H, Cui L, Tsuji N, Maeda M, Nakata K, Nogi K. Friction stir welding ofcarbon steel. Mater Sci Eng A 2006;429:50–7.

[17] Chung YD, Fujii H, Nakata K, Nogi K. Friction stir welding of high carbon toolsteel (SK85) below eutectoid temperature. Trans JWRI 2009;38:37–41.

[18] Cui L, Fujii H, Tsuji N, Nogi K. Friction stir welding of a high carbon steel.Scripta Mater 2007;56:637–40.

[19] Chung YD, Fujii H, Ueji R, Tsuji N. Friction stir welding of high carbon steelwith excellent toughness and ductility. Scripta Mater 2010;63:223–6.

[20] Ueji R, Fujii H, Cui L, Nishioka A, Kunishige K, Nogi K. Friction stir welding ofultrafine grained plain low-carbon steel formed by the martensite process.Mater Sci Eng A 2006;423:324–30.

[21] Voigt RC, Loper Jr CR. A study of heat-affected zone structures in ductile castiron. Weld J 1983:82–8.

[22] El-Banna EM. Effect of preheat on welding of ductile cast iron. Mater Lett1999;41:20–6.

[23] Fujii H, Yamaguchi Y, Kiguchi S, Nogi K. Surface hardening of cast irons byfriction stir processing. Mater Trans 2008;49:2837–43.

[24] Imagawa K, Fujii H, Morisada Y, Yamaguchi Y, Kiguchi S. Surface hardening offerritic spheroidal graphite cast iron by friction stir processing. Mater Trans2012;53:1456–60.

[25] Imagawa K, Fujii H, Morisada Y, Yamaguchi Y, Kiguchi S. Effects of toolgeometry on hardened layer of friction stir processed cast iron. Mater Trans2012;53:1952–5.

[26] Cheng TW, Lui TS, Chen LH. Microstructural features and erosion wearresistance of friction stir surface hardened spheroidal graphite cast iron. MaterTrans 2012;53:167–72.

[27] Shi J, Savas MA, Smith RW. Plastic deformation of a model material containingsoft spheroidal inclusion: spheroidal graphite cast iron. J Mater ProcessTechnol 2003;133:297–303.

[28] Krauss G. Martensite in steel: strength and structure. Mater Sci Eng A1999;273–275:40–57.

[29] Rao PP, Putatunda SK. Influence of microstructure on fracture toughness ofaustempered ductile iron. Metall Mater Trans A 1997;28:1457–70.

[30] Rao PP, Putatunda SK. Investigations on fracture toughness of austemperedductile irons austenitized at different temperatures. Mater Sci Eng A2003;349:136–49.

[31] Eric O, Sidjanin L, Miskovic Z, Zec S, Jovanovic MT. Microstructure andtoughness of CuNiMo austempered ductile iron. Marter Lett 2004;58:2707–11.