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\title{Air--Cooled Bainitic Steels for Strong, Seamless Pipes\\
Part 1: Alloy Design, Kinetics and Microstructure}
\author{G. Gomez, T. P\'erez and H. K. D. H. Bhadeshia$^{\dag}$}
\vskip 0.2cm
\begin{document}
\maketitle

\centerline{Department of Metallurgy, Tenaris Argentina R \&\ D}
\centerline{Simini 250, 2804 Campana, Pcia. Buenos Aires, Argentina}\vskip 0.2cm
\centerline{$^\dag$Materials Science and Metallurgy, University of Cambridge}
\centerline{Cambridge CB2 3QZ, U.K., www.msm.cam.ac.uk/phase--trans}

\section*{Abstract}

Seamless pipes are usually quenched and tempered after manufacture in order to obtain the desired properties. The possibility of avoiding this heat treatment is explored here, by designing steels which transform into an carbide--free bainitic microstructure during cooling from the hot processing phase of the pipe. The cooling conditions studied are appropriate for the manufacturing process and care has been taken to ensure that the correct microstructure is generated for the range of cooling rates experienced in industrial practice. The alloy design has been validated by kinetic and microstructural studies.

\section{Introduction}

Seamless pipes for applications in the oil and gas industries are made from steel in the quenched and tempered martensitic condition. This heat treatment adds to cost and it is difficult to ensure a uniform microstructure when making pipes with thick walls \cite{Gorni:2005}. An alternative scenario which avoids both of these difficulties would be to allow a carbide--free bainitic microstructure to evolve directly after the production of the seamless pipe as it cools to ambient temperature. With careful design, such microstructures have previously been shown to result in a good combination of strength and toughness both in steels \cite{Sandvik:1981,Bhadeshiaedmonds:1983a,Bhadeshiaedmonds:1983b,Miihkinen:1987a,Miihkinen:1987b,Miihkinen:1987c,Wenyan:1997,Barbacki:1998,Wei:2004,Fang:2003b,Caballero:2006b} and in cast irons \cite{Dorazil:1982,Moore:1985,Franetovic:1987}.

The purpose of the work presented here was to see whether this microstructure could be made by continuous cooling transformation in conditions appropriate for the manufacture of seamless pipes. The property requirements are a proof strength in excess of 1000~MPa with a ductile--brittle transition temperature below $-20^\circ$C (with a minimum Charpy value of 21~J and a minimum average of 24~J).
After hot rolling, the pipes are currently heat--treated by quenching to martensite and tempering. It clearly would be an advantage to eliminate the need for this heat--treatment.  As will be seen later, the properties required are demanding.  Promising alloys have been developed in previous work and in another application concept, in which the mixture of bainitic ferrite and carbon--enriched retained austenite was generated during cooling,to give yield strength greater than 1100~MPa and a fracture toughness in excess of 160~MPa\,m$^{\frac{1}{2}}$ \cite{Caballero:part1,Caballero:part2}. The cooling conditions, sample dimensions and fabrication procedures for those alloys differ from that during pipe manufacture but the previous work provides a foundation from which to proceed. 

Part 1 of the study deals with alloy design, transformation kinetics and microstructure, Part 2 with the mechanical properties achieved.

\section{Alloy Design}

The alloy design was aimed at producing a microstructure composed of bainitic ferrite and films of retained austenite during air cooling from the austenite phase field.  Finite element thermal calculations indicated that for  tube thicknesses between 8--24\,mm, the average cooling rate in air on exiting the hot rolling mill (finish rolling temperature: 950--1000$^\circ$C) is 0.1-0.5$^\circ$C\,s$^{-1}$. It was judged that to achieve a fine microstructure requires a bainite--start temperature ($B_S$) less than 500$^\circ$C, which should also help avoid products such as allotriomorphic ferrite, pearlite and Widmanst\"atten ferrite at the expected cooling rates.

 The design criteria for the steels include the following:
\begin{enumerate}
\item Sufficient silicon or aluminium addition to avoid the precipitation of cementite during the bainite reaction, thus allowing the carbon to be retained within the residual austenite \cite{Deliry:1965,Pomey:1966,Schissler:1975,Kozeschnik:2008};
\item the extent of transformation should be such that blocky regions of austenite are avoided in the final microstructure \cite{Bhadeshiaedmonds:1983a,Bhadeshiaedmonds:1983b}; this can be achieved by ensuring that the carbon concentration that the austenite can tolerate $(x_{T_0})$ before diffusionless transformation becomes impossible is at a maximum. Substitutional solute concentrations can be adjusted to maximise  $x_{T_0}$  as described elsewhere \cite{Bhadeshiaedmonds:1983a,Bhadeshiaedmonds:1983b,Caballero:part1,Caballero:part2};

\item the microstructure should be as fine as possible in order to achieve the required strength. This is usually achieved by suppressing the transformation temperature, increasing the strength of the austenite and enhancing the driving force for transformation \cite{Pickering:1958,Pickering:1992b,Singh:1998c}. Toughness is also enhanced by refining the microstructure because the transformation of carbon--enriched austenite into coarse, untempered martensite can otherwise lead to spontaneous cracking \cite{Chatterjee:cracks}. Notice that although large quantities of carbon can be added to suppress the transformation temperature and obtain nanocrystalline bainite, the strength would then become excessive for the seamless--pipe application; furthermore, the transformation times required would make this particular application impractical \cite{Caballero:current}. 
\end{enumerate}

The goal of the present work was to obtain a proof strength in excess of 1000~MPa whilst meeting the pipe Charpy toughness requirements, and for the processing conditions outlined above. Thus, the nickel--containing steels in \cite{Caballero:part1} formed the starting point of the design process, Table~\ref{tab:composition}. Time--temperature--transformation (TTT) diagrams were calculated as in \cite{Bhadeshia:1982,Bhadeshia:MS1,Bhadeshia:MS2} and these were converted into continuous cooling transformation (CCT) diagrams assuming the Schiel additivity principle \cite{Christian:2003b}.  Two compositions were selected for manufacture, Table~\ref{tab:composition}; the significant difference between these is the carbon concentration. The concentration of manganese was kept to a practical low limit (0.1~wt\%) in order to avoid banding due to solidification induced chemical segregation; such banding can have a detrimental effect on the mechanical properties since the Mn--enriched regions remain martensitic on cooling \cite{Khan:1990}. Fig.~\ref{fig:banding} shows the such banding has successfully been avoided. Nickel, on the other hand, helps improve the toughness of ferrite \cite{Jolley:1968,Norstrom:1979}. The chromium concentration was reduced in order to maximise the amount of bainitic transformation and aluminium was permitted in small concentrations for the purposes of steel deoxidation. The molybdenum addition is there to combat austenite grain boundary embrittlement due to impurities such as phosphorus \cite{Briant:1978,Olefjord:1978}. 

 
\begin{figure}[htdp]
\centering
\subfigure(a){
\includegraphics[width=7.5cm]{alloy1_opt.eps}}
\subfigure(b){
\includegraphics[width=7.5cm]{alloy2_opt.eps}}
\caption{Optical micrographs from the hot--rolled states of (a) alloy~1 and (b) alloy~2, showing the absence of microstructural banding..}
\label{fig:banding}
\end{figure}

\begin{table}[htdp]
\caption{Alloy compositions, wt\%, austenite grain size measured as a mean lineal intercept, $\overline L_\gamma$, in $\mu$m, and volume fraction $V_\gamma$ of retained austenite. In the case of the Gleeble samples, the cooling rates associated with each $V_\gamma$ value are also stated. The data for the as--rolled samples will be discussed in Part~2 \cite{Gomez:2008b}.}
\begin{center}
\begin{tabular}{lccccc}\midrule
& Previous \cite{Caballero:part1} & \multicolumn{2}{c}{Alloy 1} &  \multicolumn{2}{c}{Alloy 2} \\
& & Specified & Actual & Specified & Actual \\\midrule
C & 0.30 & 0.25 & 0.24 & 0.30 & 0.30 \\[3pt]
Mn & --& 0.1 & 0.09 & 0.1 & 0.10 \\[3pt]
Si & 1.5 & 1.40 & 1.27 & 1.40 & 1.42 \\[3pt]
Cr & 1.44 & 1.00 & 1.00 & 1.00 & 1.03 \\[3pt]
Mo & 0.25& 0.25 & 0.23 & 0.25 & 0.22 \\[3pt]
Ni &3.5 &  3.60 & 3.64 & 3.60 & 3.48 \\[3pt]
Al & & 0.03 & 0.014 & 0.03 & 0.040 \\[3pt]
S& &  $<0.002$ & 0.001 & $<0.002$ & 0.002 \\[3pt]
P & & $<0.01$ & 0.005 & $<0.01$ & 0.01 \\[3pt]
N & & $<0.006 $ & 0.0063 & $<0.006$ & 0.0023 \\[3pt]
O & & $<0.003$ & 0.0014 & $<0.003 $ & 0.0011 \\[10pt]
$\overline L_\gamma$, as--rolled  & & \multicolumn{2}{c}{$52\pm 10$} & \multicolumn{2}{c}{$39\pm 11$}\\[3pt]
$\overline L_\gamma$, normalised & & \multicolumn{2}{c}{$22\pm 3$} & \multicolumn{2}{c}{$16\pm 4$}\\[3pt]
$\overline L_\gamma$, Gleeble & & \multicolumn{2}{c}{$47\pm 12$} & \multicolumn{2}{c}{$36\pm 10$}\\[10pt]
$V_\gamma$, as--rolled & & \multicolumn{2}{c}{$0.18\pm 0.01$} &\multicolumn{2}{c}{$0.13\pm 0.01$} \\[3pt]
$V_\gamma$, Gleeble, 0.1${\rm ^\circ C\,s^{-1}}$ & & \multicolumn{2}{c}{$0.28\pm 0.03$} &\multicolumn{2}{c}{$0.27\pm 0.03$} \\[3pt]
$V_\gamma$, Gleeble, 0.2${\rm ^\circ C\,s^{-1}}$ & & \multicolumn{2}{c}{$0.18\pm 0.03$} &\multicolumn{2}{c}{$0.15\pm 0.03$} \\[3pt]
$V_\gamma$, Gleeble, 0.5${\rm ^\circ C\,s^{-1}}$ & & \multicolumn{2}{c}{$0.10\pm 0.02$} &\multicolumn{2}{c}{$0.05\pm 0.02$} \\[3pt]
$V_\gamma$, Gleeble, 2.0${\rm ^\circ C\,s^{-1}}$ & & \multicolumn{2}{c}{--} &\multicolumn{2}{c}{--} \\[3pt]

\midrule  
\end{tabular}
\end{center}
\label{tab:composition}
\end{table}

Figure~\ref{fig:CCT} shows the calculated transformation--start temperatures as a function of the cooling rate. It is  therefore expected that both steels would present predominantly bainitic microstructures during cooling at 0.1--0.5$^\circ$C\,s$^{-1}$. Some allotriomorphic ferrite is possible at the lower end of this cooling range (0.1-0.2$^\circ$C\,s$^{-1}$) but this is expected to be minimal given the high hardenability of the steel.  The calculated $B_S$ temperatures were below 500$^\circ$C in both cases: 471$^\circ$C   and 446$^\circ$C for Alloys 1 and 2 respectively. The effect of carbon is to suppress transformation temperatures at all cooling rates, so a finer microstructure is expected in Alloy~2.
 

\begin{figure}[!ht]
  \begin{center}
    \begin{minipage}[t]{0.72\linewidth}
\raisebox{-6cm}{\includegraphics[width=0.98\linewidth]{CCT.eps}} 
    \end{minipage}\hfill
    \begin{minipage}[t]{0.25\linewidth}
      \caption{Calculated transformation start temperatures for alloys 1 and 2. $\alpha$, $\alpha_B$ and $\alpha'$ stand for allotriomorphic ferrite, bainite and martensite respectively.}  \label{fig:CCT}
    \end{minipage}
  \end{center}  
\end{figure}


\section{Experimental Details}

The alloys were prepared as 20~kg melts (Table~\ref{tab:composition}) in a vacuum induction furnace. The materials were electroslag remelted in order to reduce the sulphur content.  Slabs of 140 mm thickness were cropped to 123~mm thickness due to limitations in the pilot mill and then hot rolled into plates of 16~mm thickness. During rolling the reheating and finish rolling temperatures were chosen to be 1250$^\circ$C and 1000$^\circ$C to approximately match those used during the industrial processing of pipes. After hot rolling the plates with dimensions $800 \times 170 \times 16$\, mm, were air cooled to room temperature.

Hardness, tensile and impact tests were conducted on the as--rolled samples. Cylindrical samples of 8.9~mm diameter  and 35~mm of guage length were used for tensile testing, with the results in each case averaged over two samples. Charpy V-notch tests were conducted on full size samples $(10\times10\times55$~mm) machined in the transverse direction (``T-L" according to ASTM E23). The reported impact results correspond to average values over 3 measurements for each temperature, which ranged from room temperature to -40$^\circ$C. The percentage of ductile fracture on the broken Charpy specimens was measured according to standard practice \cite{ASTME23} following industrial specifications. Vickers hardness was measured using a 20~kg load, and averaged over 5 measurements at random positions.

Optical and scanning electron microscopy studies were on cold mounted, ground and polished samples and etched in 2\% nital. The amount of retained austenite in the hot--rolled samples \cite{Gomez:2008b} was determined using quantitative X-ray analysis on a Phillips X'Pert diffractometer and unfiltered Cu $K_\alpha$ radiation. The scanning speed was 0.1$^\circ$C\,min$^{-1}$ over the range $2\theta = 30-110^\circ$. The machine was operated at 40~kV and 40 mA. The retained austenite content was calculated from the integrated intensities of (111), (002), (022), (113) and (222) austenite peaks, and those of (011), (002), (112) and (022) planes of ferrite, using Rietveld analysis. The reheated Gleeble samples were, because of their small size, analysed using just two peaks (111)$_\gamma$ and (011)$_\alpha$ by scanning over $2\theta = 40-55^\circ$ at the same speed, so in this case the results are less accurate but should indicate any trends correctly.

Continuous cooling transformation (CCT) diagrams were determined from dilatometric tests performed in a Gleeble 3500 thermomechanical simulator, using 6~mm diameter cylindrical samples heated at 5$^\circ$C\,s$^{-1}$ up to 1000$^\circ$C (chosen to match the austenite grain size to the rolled plates), followed immediately by  cooling at a fixed rate down to room temperature.  The tests were conducted in a vacuum--purged chamber under flowing high purity argon containing 5\% hydrogen to prevent decarburisation and oxidation. The temperature evolution was followed with a type $K$--type thermocouple spot--welded on the surface of the specimen, in its middle section. The changes in sample diameter were acquired as a function of time with a high resolution dilatometer at the location of the thermocouple (the accuracy of the measurement was $\pm$0.1\,$\mu$m). Cooling rates in in the range 0.1--5$^\circ$C were used to construct CCT diagrams and the fraction of transformation was determined by applying the lever rule to the dilatometer reading relative to the austenite and ferrite lines from regions where transformation was absent. The Vickers hardness was measured using a 10\,kg load, and averaged over at least 3 measurements at random positions in each case.

The tempering behaviour was studied on samples $20\times 20 \times 8\,$mm, sealed in quartz tubes containing a partial pressure of argon,  at temperatures between 250--500$^\circ$C for 1\,h.  Further  heat treatments were carried out on $200\times 170\times 16\,$mm plates which were then machined into tensile and impact test samples: 
\begin{itemize}
\item Alloy 1, 880$^\circ$C for 15\,min followed by air cooling (normalising).
\item Alloy 2, 840$^\circ$C for 15\,min followed by air cooling (normalising).
\item Alloy 2, tempering at 300$^\circ$C for 1\,h, followed by air cooling.
\item Alloy 2, tempering at 500$^\circ$C for 30\,min followed by air cooling.
\end{itemize}
Transmission electron microscopy was conducted on a JEOL 2000FX transmission electron microscope operating at 200\,kV. Thin foil preparation methods have been described elsewhere \cite{Bhadeshia:1979c}.


\section{CCT diagrams}

Examples of the raw data obtained from Gleeble tests are presented Fig.~\ref{fig:dilatometer}, for two cooling rates. Austenite formation began on heating at $Ac_1 = 725~\hbox{and}~735^\circ$C for Alloys~1 and 2 respectively,  and finished at $Ac_3 = 835-845^\circ$C for alloy 1 and or 820--830$^\circ$C for Alloy~2. The maximum reheating temperature of 1000$^\circ$C is far above $Ac_3$ leading  to relatively coarse austenite grains.  The measured austenite grain sizes are given in Table~\ref{tab:composition}; it is important to note that the Gleeble samples have essentially the same $\overline L_\gamma$ as the as--rolled specimens so that the CCT diagrams are representative of the rolled condition. The reason why $\overline L_\gamma$ for Alloy~2 is consistently smaller than that for Alloy~1 is not clear, but it is speculated that the reduction is associated with its higher aluminium concentration which via the formation of aluminium nitride, acts to refine the grain size \cite{Wilson:1988,Walker:1990,Gladman:1994,Eldridge:1998}.
  
\begin{figure}
\centering
\subfigure(a){
\includegraphics[width=7.5cm]{dilatometer1.eps}}
\subfigure(b){
\includegraphics[width=7.5cm]{dilatometer2.eps}}\\
\subfigure(c){
\includegraphics[width=7.5cm]{dilatometer3.eps}}
\subfigure(d){
\includegraphics[width=7.5cm]{dilatometer4.eps}}
\caption{Typical dilatometric data. (a,b) Alloy 1; (c,d) Alloy 2.}
\label{fig:dilatometer}
\end{figure}

The dilatometric measurements following austenitisation at 1000$^\circ$C in combination with microscopy were used to measure the CCT diagrams illustrated in Fig.~\ref{fig:measuredCCT}. The temperatures corresponding to  5\%, 10\%, 50\% and 90\% of transformation are plotted as a function of the cooling time from 830$^\circ$C (the approximate temperature at which phase transformation can begin). In some cases a surface decarburised layer some 70\,$\mu$m thickness occurred, which represented about 0.03 volume fraction of the sample; this is why at 1\%\ transformation temperature is not represented in  Fig.~\ref{fig:measuredCCT}.

The $M_S$ temperatures were determined from the change in the slope of the transformation strain versus temperature curve during continuous cooling. $M_S$ is not plotted for Alloy~1 because fully martensitic transformation was not obtained for the cooling rates studied; transformation during cooling at 5$^\circ$C\,s$^{-1}$ began at about 500$^\circ$C although in excess of 95\% of the reaction occurred below 350$^\circ$C when an abrupt rise in decomposition rate of austenite was noted in the dilatometric curve. This is consistent with the $M_S$ temperature in the range 350--355$^\circ$C calculated as in \cite{Bhadeshia:MS1,Bhadeshia:MS2}  and using the Andrews equation \cite{Andrews:1965}. The extent of transformation above the calculated $M_S$ increased as the cooling rate decreased and typical microstructures for the slow cooling rates of relevance are illustrated in Fig.~\ref{fig:micalloy1}. At 0.5$^\circ$C\,s$^{-1}$ most of the transformation took place above the $M_S$ and below the calculated $B_S$ (471$^\circ$C). The same occurred when cooling at  0.2$^\circ$C\,s$^{-1}$, as is clear from the dilatometric curve in Fig.~\ref{fig:dilatometer}. The appearance of bainite was confirmed metallographically, Fig.~\ref{fig:micalloy1}. Allotriomorphic ferrite observed at 0.2$^\circ$C\,s$^{-1}$ and 0.1$^\circ$C\,s$^{-1}$ but  the fraction was less than 0.05 as estimated using point counting \cite{DeHoff:1968}.

The retained austenite contents reported in Table~\ref{tab:composition} are generally consistent with the dilatometric analysis; $V_\gamma$ decreases as the cooling rate is increased because bainite then has less time to form, thus resulting in a greater tendency to form martensite. Alloy~2 in general has a smaller fraction of austenite, consistent with its greater hardenability (higher average carbon concentration). 

\begin{figure}
\centering
\subfigure(a){
\includegraphics[width=7.5cm]{Alloy1.eps}}
\subfigure(b){
\includegraphics[width=7.5cm]{Alloy2.eps}}
\caption{Measured CCT diagrams. (a) Alloy~1, austenitised at 1000$^\circ$C; (b) Alloy~2 austenitised at 1000$^\circ$C.}
\label{fig:measuredCCT}
\end{figure}


\begin{figure}
\centering
\subfigure(a){
\includegraphics[width=7cm]{Alloy1_1.eps}}
\subfigure(b){
\includegraphics[width=7cm]{Alloy1_3.eps}}\\
\subfigure(c){
\includegraphics[width=7cm]{Alloy1_2.eps}}
\subfigure(d){
\includegraphics[width=7cm]{Alloy1_4.eps}}
\caption{Alloy 1. (a) 1$^\circ$C\,s$^{-1}$; (b) 0.5$^\circ$C\,s$^{-1}$; (c) 0.2$^\circ$C\,s$^{-1}$; (d) 0.1$^\circ$C\,s$^{-1}$.}
\label{fig:micalloy1}
\end{figure}

\begin{figure}
\centering
\subfigure(a){
\includegraphics[width=7cm]{Alloy2_1.eps}}
\subfigure(b){
\includegraphics[width=7cm]{Alloy2_2.eps}}\\
\subfigure(c){
\includegraphics[width=7cm]{Alloy2_3.eps}}
\subfigure(d){
\includegraphics[width=7cm]{Alloy2_4.eps}}
\caption{Alloy 2. (a) 1$^\circ$C\,s$^{-1}$; (b) 0.5$^\circ$C\,s$^{-1}$; (c) 0.2$^\circ$C\,s$^{-1}$; (d) 0.1$^\circ$C\,s$^{-1}$.}
\label{fig:micalloy2}
\end{figure}
 
The general dependence of Alloy~2 on the cooling rate was similar to that of Alloy~1 although with its higher carbon concentration, most of the austenite decomposition occurred below 320$^\circ$C for $\dot T>2^\circ$C\,s$^{-1}$ with fully martensitic microstructure, although a tiny amount of bainite could be seen at the decarburised surface.  Fine bainite was observed for $\dot T<2^\circ$C\,s$^{-1}$ (Fig.~\ref{fig:micalloy2}) but the microstructure also contained some martensite as inferred from the dilatometric curves, Fig.~\ref{fig:dilatometer}. Another difference with the lower carbon alloy was the degree of microstructural refinement: when comparing micrographs in Figs.~\ref{fig:micalloy1} and \ref{fig:micalloy2}, it is clear that finer bainite was obtained with the higher carbon content.

There is a refinement of microstructure in Alloy~2, when compared with Alloy~1, through the suppression of transformation to lower temperatures; this is evident by comparing the 50\%\ reaction CCT curves for the two steels (Fig.~\ref{fig:measuredCCT}).  The difference is more obvious in the dilatometric plots in Fig.~\ref{fig:dilat2}; the transformation begun at nearly the same temperature in both steels, but its development was suppressed in Alloy~2 by about 50$^\circ$C. 
 


\begin{figure}[!ht]
  \begin{center}
    \begin{minipage}[t]{0.72\linewidth}
\raisebox{-4cm}{\includegraphics[width=0.98\linewidth]{dilat2.eps}} 
    \end{minipage}\hfill
    \begin{minipage}[t]{0.25\linewidth}
      \caption{Dilatometric curves during cooling from 1000$^\circ$C at $2^\circ$C\,s$^{-1}$.}  \label{fig:dilat2}
    \end{minipage}
  \end{center}  
\end{figure}

In summary, there is reasonable agreement between the calculated and measured CCT curves, notwithstanding the small amount of transformation associated with the decarburised layers in the dilatometric samples. Table~\ref{tab:comparison} contains a further comparison, where $\dot T_\alpha$ is the minimum cooling rate to avoid allotriomorphic ferrite, and $\dot T_{\alpha'}$ the minimum cooling rate needed to achieve a fully martensitic microstructure. It is particularly useful that the range of cooling rates in which bainite is the predominant transformation is correctly estimated. 

\begin{table}[htdp]
\caption{Comparison of calculations and measured data from CCT diagrams.}
\begin{center}
\begin{tabular}{ccccccccc}\midrule
Steel & 	\multicolumn{2}{c}{$\dot T_\alpha$ / $^\circ$C\,s$^{-1}$}	& \multicolumn{2}{c}{$\dot T_{\alpha'}$ / $^\circ$C\,s$^{-1}$}&	\multicolumn{2}{c}{$M_S$ / $^\circ$C}& \multicolumn{2}{c}{$B_S$ / $^\circ$C}\\[3pt]
	&Calc. &	Meas. &	Calc.& Meas. &	Calc.	& Meas. &	Calc. & Meas. \\\midrule
Alloy 1 & 	0.25 & 	0.2-0.5& 	8& 	$> 5$&	355--350&	350&	471&	500\\[6pt]
Alloy 2	&0.22&	0.2--0.5&	6	&2--5	&330--340&	320& 	446& 	500\\[3pt]\midrule
\end{tabular}
\end{center}
\label{tab:comparison}
\end{table}%

Fig.~\ref{fig:hardness} shows the hardness data as a function of the alloy and cooling rate. The hardness expected from a fully martensitic microstructure is also plotted as a function of the cooling rate using an equation due to Maynier {\em et al.} \cite{Maynier:1978} which accounts for the chemical composition of the steel and $\dot T$. The results confirm that the microstructure is essentially martensitic when $\dot T\geq 1^\circ$C\,s$^{-1}$. 


\begin{figure}[!ht]
  \begin{center}
    \begin{minipage}[t]{0.72\linewidth}
\raisebox{-6cm}{\includegraphics[width=0.98\linewidth]{hardness.eps}} 
    \end{minipage}\hfill
    \begin{minipage}[t]{0.25\linewidth}
      \caption{Hardness as a function of the cooling rate.}  \label{fig:hardness}
    \end{minipage}
  \end{center}  
\end{figure}

\section{Conclusions}

It has been demonstrated that for a range of cooling rates appropriate in the manufacture of seamless pipes, a microstructure which consists mainly of bainite can be obtained in two alloys which differ only in the carbon concentration in the range 0.25--0.30\,wt\%. The higher carbon concentration alloy is found to be more reliable in generating the required microstructure and has the additional advantage that the overall structure is finer. 

As an additional outcome, the metallographic and kinetic measurements confirm the calculated continuous cooling transformation diagrams. The mechanical properties of the alloys are explored in Part 2 of this paper \cite{Gomez:2008b}.

\section*{Acknowledgements}

The authors are grateful to Tenaris Argentina (Research and Development) for funding this work.

\begin{thebibliography}{10}
\expandafter\ifx\csname url\endcsname\relax
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 \end{document} 