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\begin{document}
\begin{frontmatter}
\title{The effects of dilution and base plate strength on stress distributions in multi-pass welds deposited using low transformation-temperature filler alloys}

\author[cambridge]{T. I. Ramjaun}\ead{tir20@cam.ac.uk}
\author[cambridge]{H. J. Stone}
\author[West]{L. Karlsson}
\author[ISIS]{J. Kelleher}
\author[cambridge]{S. W. Ooi}
\author[ESAB]{K. Dalaei}
\author[Germany]{J. Rebelo Kornmeier}
\author[cambridge]{H. K. D. H. Bhadeshia}
\address[cambridge]{Materials Science and Metallurgy,
University of Cambridge, CB3 0FS, U.K.}
\address[West]{Engineering Science, University West, SE-461 86 Trollh\"attan, Sweden}
\address[ISIS]{ISIS Facility, Rutherford Appleton Laboratory, Didcot, OX11 0QX, U.K. }
\address[ESAB]{ESAB AB, Lindholmsallen 9, 40277, Gothenburg, Sweden}
\address[Germany] {Technische Universit\"at M\"unchen, Forschungsneutronenquelle Heinz
Maier-Leibnitz (FRM II), 85747 Garching, Germany}
\date{\today}


\begin{abstract}
Transformation plasticity can be utilised to control residual stresses  in steel welds. This requires  special filler alloys that transform at a sufficiently low temperature to compensate for accumulated thermal contraction strains. However, the welding parameters needed to optimise the effect in multi-pass joints have yet to be established. This topic has been investigated by characterising the residual stress distribution in multi-pass welds fabricated with different welding alloys and base plates using neutron diffraction to assess the effects of dilution and base plate strength. Whilst the use of richly-alloyed weld metal does enhance fatigue performance in single-pass joints, the extent of stress relief that can be derived from transformation plasticity is reduced due to incomplete martensitic transformation when further layers are deposited. For all cases studied, compressive stresses were measured in the weld metal with balancing tensile stress in the heat-affected zone of the plate. The magnitude of the tension was observed to be a function of the strength of the base plate. Recommendations are also presented for the combination of welding and material parameters that lead to the optimum exploitation of transformation plasticity as a method for boosting the fatigue performance of multi-pass welded joints. 
\end{abstract}

\begin{keyword}
transformation induced plasticity \sep martensitic weld metal \sep residual stress \sep welding  \sep neutron diffraction\end{keyword}
\end{frontmatter}

\clearpage
\section{Introduction}

Residual stresses that accumulate during the cooling of welded joints are an anathema in engineering design because they limit the loads that can be applied in service and often reduce fatigue performance. Post-weld treatments can be employed to counter these thermally induced stresses but an alternative method, proposed originally by Alberry and Jones \cite{Jones:1977c}, is to exploit the strains associated with solid-state phase transformation of the weld filler to cancel the thermal contraction \cite{Ohta:1999,Eckerlid:2003,Lixing:2004, Zenitani:2007,Francis:2007,Darcis:2008,Asprino:2008,Dai:2008,Mikami:2009, Kromm:2011,Altenkirch:2011,Kasuya:2012}. This approach aims to improve the welding fabrication process through making expensive and labour intensive post-weld surface/heat treatments redundant. 

As a weld cools, it contracts, producing harmful tensile residual stresses. However, the displacive nature of the austenite to martensite phase transformation ($\gamma \to \alpha'$) is capable of negating these stresses and even producing compression in the weld metal. The details of this mechanism have been reviewed extensively elsewhere \cite{Francis:2007b, Ooi:2014} but an essential requirement of such weld filler alloys is that the transformation plasticity should not be exhausted until the weld cools to ambient temperature. This usually means that the weld metal should be designed to be martensitic with a low martensite-start temperature ($M_\mathrm{S}$). If transformation occurs at an elevated temperature, further thermal contraction will lead to a build up of tensile stress as the weld continues to cool, thus eliminating the initial stress relaxation benefits.


Naturally, other essential properties such as toughness and strength should not be compromised in the alloy design process  \cite{Shirzadi:2009b,Moat:2011} and the carbon concentration must be low enough to avoid the embrittlement associated with hard martensite. With care, it is possible to reduce residual stresses in critical locations and enhance the fatigue strength of the joint  \cite{Francis:2009,Francis:2009b,Karlsson:2010,Karlsson:2011c,Miki:2012,Takahashi:2013b,Miki:2013}. Whilst much progress has been made in this field, open questions remain about the state of the joint when more than one layer of weld metal is deposited in order to fill the gap and the consequence of the multiple heat treatments experienced by underlying weld metal. The purpose of the work presented here is to elucidate these issues through characterisation of the residual stresses produced in multi-pass welds as a function of the transformation characteristics of the weld metal and the mechanical strength of the base plate.

\section{Experimental Method}

Six 3-pass welds were prepared from three different steel plates (`355'--a low strength structural steel, Weldox 700, Weldox 960) with three different welding consumables (LTT-1, LTT-2, Coreweld 89), as shown in Table~\ref{tab:samples}. LTT-1 and LTT-2 are martensitic stainless steel filler alloys with a low $M_\mathrm{S}$, whilst Coreweld 89 is a commercially available filler with a much higher $M_\mathrm{S}$. Compositions of the plates and fillers are shown in Table~\ref{tab:compositions}. LTT-2 is highly alloyed and was designed to compensate for dilution, which inevitably occurs when the filler mixes with the base plate. The martensite-start temperature of an undiluted sample of LTT-1 was measured using dilatometry and determined to be $164\pm 12\,^\circ$C. This value is in agreement with the calculation by Steven \& Haynes and the $M_\mathrm{S}$ of the undiluted alloys are also included in Table~\ref{tab:compositions}. The mechanical properties of the welding fillers and base plates are in Table~\ref{tab:mech_prop}, the LTT data are measured values. 


Each of the welds was produced by mechanised gas-shielded metal arc welding (GMAW) on plates of dimensions $500\times150\times15$\,mm, machined with a 60$^\circ$ V-groove along the long direction and a  root radius of 4\,mm, Fig.~\ref{fig:plate}. The plates were clamped to the bench prior to welding, which was performed horizontally in the down hand position with an initial pre-heat of 50$^\circ$C and an inter-pass temperature of 100--125$^\circ$C. All three welding alloys were deposited using metal-cored electrode wire with an Ar/CO$_2$ shielding gas. The heat input for the LTT and Coreweld 89 welding alloys was $\sim$1.0 and $\sim$1.5 kJ\,mm$^{-1}$ respectively.

It is the alloying additions that control the $M_\mathrm{S}$ of the filler so four plates  were fabricated to understand the effects of weld metal dilution with the base plate. An additional two welds were prepared  to determine the effects of base plate strength with respect to the filler.


\subsection{Neutron Diffraction}

The residual stresses in the welded plates were measured using neutron diffraction \cite{Fitzpatrick:2003} on two strain scanning instruments: the STRESS-SPEC beamline \cite{Hofmann:2006} at the FRM~II reactor (weld 4, 5, 6), and the ENGIN-X beamline \cite{Santisteban:2006} at the ISIS pulsed spallation neutron source (weld 1, 2, 3). For both experiments, strain scanning was performed across a plane perpendicular to the weld at the position along the sample shown in Fig.~\ref{fig:weldcomb}a. Measurements were made at 2.5, 5, 7.5, 10 and 12.5\,mm below the top surface across both sides of the weld, heat-affected zone (HAZ) and base plate to identify any asymmetry in the stress field. It was assumed that the principal stresses are parallel to the plate edges.

The neutron diffraction techniques and subsequent stress analysis followed the same procedure as detailed elsewhere \cite{Ramjaun:2014} but the key experimental parameters are shown in Table~\ref{tab:expt_parameters}, where $\lambda$ is the neutron wavelength, 2$\theta_\mathrm{M}$ is the Bragg angle of the monochromator and $V_V$ is the gauge volume for the appropriate measurement direction.

\subsection{Stress Analysis}

Strain-free reference specimens (`comb' samples) were used to account for the effect of compositional variation between the filler alloy and base plate for the measured lattice spacings. These were produced as 3\,mm thick cross-sectional slices from the welded plate, Fig.~\ref{fig:weldcomb}a, and further slotted at 3\,mm intervals along the transverse direction using wire electro-discharge machining (EDM), Fig.~\ref{fig:weldcomb}b. Due to the position of the EDM cuts, it was not possible to achieve perfect correspondence between the measurement positions in the welded plates and the combs. Therefore, a linear interpolation was performed between the comb measurement positions at each depth to calculate the strain-free lattice spacings, $d_{0,hkl}$, for the relevant locations in the welded plates.

The elastic strain, $\varepsilon_{hkl}$, at each measurement location and direction was then found as follows:
\begin{equation}
\varepsilon_{hkl} = \frac{d_{hkl} - d_{0,hkl}} {d_{0,hkl}}
\end{equation}
where $d_{hkl}$, is the interplannar spacing in the weld for a reflection ($hkl$). Elastic strains were calculated from the ENGIN-X data in an analogous way using the lattice parameters.

The stress, $\sigma_{ii}$, in each of the three orthogonal directions was found from the strain measurements according to:
\begin{equation}
\sigma_{ii} = \frac{E} {(1 + \nu)} \left[ \varepsilon_{ii} + \frac{\nu} {(1 - 2\nu)} (\varepsilon_{11} + \varepsilon_{22} + \varepsilon_{33})\right] \label{eqn:stress}
\end{equation}
where $E$ is the Young's modulus and $\nu$ is Poisson's ratio. $E=220$\,GPa and $\nu=0.28$ are the diffraction elastic constants for the \{211\} used to calculate the stress in welds 4, 5 and 6 \cite{Eigenmann:1996}.  $E=212.7$\,GPa and $\nu=0.30$ are the macroscopic bulk properties for a ferritic steel following the Kr\"oner model and used to calculate the stress in welds 1, 2 and 3 \cite{Hutchings:2005}. $i = 1,\,2,\,3$  denotes the direction of the initial lattice parameter of lattice spacing measurement and hence, strain and stress direction with respect to the welded plate geometry.
 

\section{Results and Discussion}

\subsection{Residual Stress Measurements}

Residual stresses were measured in the longitudinal, transverse and normal directions for each of the six welded plates, Fig.~\ref{fig:contour}. The original measurement locations are indicated with crosses and the strains measured at these positions were used to map the stresses across the weld bead, HAZ and into the base plate at depths of 2.5--12.5\,mm below the surface. Analogous to previous studies \cite{Francis:2007}, the stress contours show symmetry across the centreline and follow the geometry of the weld.

Compressive stresses are evident in the weld material for all five specimens in the longitudinal direction that were welded with an LTT alloy. This is in contrast to weld 4 and confirms the shear and dilatational strains associated with the $\gamma \to \alpha'$ transformation \cite{Francis:2007} at low temperature have compensated for the tensile thermal contraction strains, which would otherwise evolve. The stresses developed in the transverse and normal directions show evidence of compression in the LTT weld material but they are much lower in comparison. The longitudinally generated stresses are greatest because this is the direction of maximum thermal constraint during cooling. Weld 4 also appears to accumulate only minor compressive stresses in the weld bead for the transverse and normal directions, although this is not necessarily an indication of low temperature transformation.

With respect to the longitudinal direction, the five LTT welds have a common stress distribution with compression in the weld metal nearer the surface and peak tensile stresses in the HAZ both below and either side of the weld. The tensile stresses are generated as a result of stress equilibration and the absence of low temperature $\gamma \to \alpha'$ transformation in these locations. The re-austenitised regions in the base plate that surround the fusion boundary will transform to martensite at an elevated temperature, thus accumulating tensile residual stresses during the subsequent thermal contraction on cooling.  Weld 4 shows a similar trend to the LTT welded plates but does not develop compressive stresses in the weld metal.

The stress contours are complemented by longitudinal-stress profiles, Fig.~\ref{fig:stress_plots}, which show the residual stresses at exact locations and are discussed in more detail in the sections that follow. The welds have been collated into two series, assessing the effects of dilution and base plate strength on residual stress with the view to improving fatigue performance. It is for this reason that particular attention is paid to the stresses nearest the specimen surface at the fusion boundary, where fatigue failure predominantly occurs. An additional figure is presented, Fig.~\ref{fig:extra}, which investigates the issue of steep stress gradients across the fusion boundary/HAZ and the number of measurement positions necessary to adequately describe the stress change. 

\subsection{The Effects of Dilution}

Welds 4, 1 and 2 were welded with increasingly-alloyed filler metals and hence, decreasing $M_\mathrm{S}$. Weld~3 comprised of three layers with only the final, capping pass being an LTT alloy, to ascertain the possibility of influencing the surface residual stresses whilst minimising the use of these relatively expensive materials. Although the stress state is evaluated throughout the entire weld in three orthogonal directions, emphasis is placed on the longitudinal stresses nearest the top surface due to these stresses being the greatest in magnitude and the surface region being the predominant initiation site for fatigue failure. 

Weld 4 was fabricated with a conventional high $M_\mathrm{S}$ filler alloy and shows the expected tensile stresses in the fusion zone that are generated because the $\gamma \to \alpha'$ transformation occurs at a relatively high temperature, thus allowing further thermal contraction strains to develop after the transformation is exhausted. A large region of high tensile stress (+600\,MPa) envelops the majority of the HAZ, extending to the measurement position nearest the top surface. This is in contrast to welds 1 and 2, which generate compression in the weld metal and have much smaller regions of tensile stress above 600\,MPa that are situated further from the top surface. Weld 3 also develops compressive stress in the final LTT alloy capping pass. 

Due to the tensile stresses developed in the weld metal for weld~4 and the  compression (-200\,MPa) in the same region for weld 1 (low $M_\mathrm{S}$ filler), it might be expected that weld~2, which has an even lower $M_\mathrm{S}$, exhibit greater compressive stresses. The data for welds 4, 1 and 2 (Fig.~\ref{fig:CFGH_profile}) are inconsistent with this hypothesis because the magnitude of compressive residual stress in weld~1 is greater than weld~2. This is caused by untransformed austenite in weld 2 resulting from the heavily alloyed LTT-2 filler with a  suppressed $M_\mathrm{S}$ \cite{Wang:2002c}. The full extent of transformation plasticity clearly cannot be achieved with the partial transformation of austenite so only minor compressive stresses are produced in weld~2. 

In order to illustrate this point, {Fig.~\ref{fig:stress_temp} provides a schematic diagram of the effect of $M_\mathrm{S}$ on the residual stresses produced in  three different welding alloys during a constrained cooling situation. A bar of each alloy is cooled from a stress-free state at high temperature with the ends of the bar fixed, preventing any thermal shrinkage. This is an idealised model of the constrained cooling that occurs in the weld and HAZ during welding \cite{Satoh:1970}. The welding filler with a high $M_\mathrm{S}$ transforms prematurely, which leads to tensile residual stress in the weld metal as it continues to thermally contract until it reaches ambient temperature. The filler with an overly suppressed $M_\mathrm{S}$, such that the $\gamma \to \alpha'$ transformation is largely incomplete at ambient temperature, achieves only partial stress relaxation. Although, this is preferable to tensile residual stress. The `ideal' filler transforms at the optimum temperature to allow complete transformation, which results in maximum stress alleviation to the extent of inducing large compressive stress in the weld metal.

The macrograph in Fig.~\ref{fig:Fmacro} shows evidence of continuous columnar grains growing across the fusion boundary between welding passes 2--3, which would indicate that the $M_\mathrm{S}$ of the weld bead in this region is similar to the inter-pass temperature \cite{Ramjaun:2014}. The welding alloy developed by Karlsson \it{et al.} \rm \cite{Karlsson:2011c} to compensate for dilution with the base plate displays greater fatigue performance than the standard LTT alloy. Importantly, this is for single-pass welds where the extent of dilution is typically 30\% and elevates the $M_\mathrm{S}$. However, LTT-2 may be not suitable as the sole filler for multi-pass welds since its undiluted $M_\mathrm{S}$ is too low for complete transformation. 

Weld 3, which incorporated a single capping pass of LTT-2, shows the desired accumulation of compressive stresses in the LTT weld metal but the remainder of the stress distribution is synonymous with weld 4. Tensile stresses are evident around the HAZ at the measurement positions nearest the surface and Fig.~\ref{fig:CFGH_profile} suggests that the maximum tensile stress for weld 3 is the same as that for weld 4. Although the final LTT-2 capping pass is unable to reverse the effects of accumulated tensile stress in the initially deposited layers of the high $M_\mathrm{S}$ filler and surrounding HAZ, compressive stresses were developed in the LTT weld metal. Similarities between the maximum tensile stress and distributions of weld 3 and 4 are evident for this work, however, the residual stress states at the actual surface could deviate.

Weld 1 and 2 appear to have the most favourable stress distributions (compared with weld 4 and 3) based on smaller regions of high tensile stress further below the top surface. Despite the benefits to fatigue performance with a single-pass, multi-pass welds produced with the highly alloyed LTT-2 may not be appropriate because of the mechanical properties associated with high levels of retained austenite. In addition, results from an investigation into the effects of inter-pass temperature ($T_\mathrm{I}$) \cite{Ramjaun:2014} for the filler LTT-1 show a further improvement of stress distribution when the $T_\mathrm{I}>M_\mathrm{S}$. Depending on the welding speed and number/geometry of passes, it is possible that the entire weld will subsequently transform simultaneously on cooling, thus realising the full benefits of transformation plasticity across the entire weld.

\subsection{Effects of Base Plate Strength}

In order to determine the versatility of the LTT alloys with regards to welding steel plates of varying strength, two base materials were selected that had yield strengths above (weld 5) and below (weld 6) that of the filler alloy. For weld 1, the yield strength of the filler was comparable to the base plate and the mechanical properties are listed in Table~\ref{tab:mech_prop}. All of the welds show compression in the weld metal and maximum tensile stresses in the vicinity of the HAZ that correlate with the yield strength of the associated base plate. This would suggest that the thermal stresses generated in the constrained weld metal and surrounding base plate approach or exceed the yield strengths of the respective materials, although there is likely to be  some difference in microstructure/properties between the HAZ and unaffected base plate. Remarkably, the phase transformation in the LTT weld metal is sufficient to negate the tensile stresses that develop in the fusion zone. However, the accumulated tensile stresses in the base plate remain, although expansion of the weld metal will inevitably lead to some alleviation and re-distribution.

Comparing the stress distribution in the HAZ for welds 1 and 4, both deposited on Weldox 700, there is a significant reduction in the magnitude of tensile stress when the LTT welding alloy is used. Applying this principle to weld 5, which has a small region of tensile stress in excess of 800\,MPa, it is possible that the entire HAZ will be above this value when welding with a conventional filler. 


\subsection{Strain Measurement Positions}

As a result of the time limitations associated with strain scanning using neutron diffraction, a compromise usually exists between the number of measurement positions and the number of specimens that can be analysed. Additional strain measurements were taken for weld 5, Fig.~\ref{fig:extra}, to ensure that the initial concentration of points was sufficient to adequately capture the stress profiles across  regions of abrupt change, such as the fusion boundary and HAZ. For the collection of welds in this work, readings were taken every 3\,mm, signified by the solid markers. Additional points have been included (open markers) to show the change in profile for measurement positions every 1\,mm. The initial concentration of points provides a satisfactory representation of the stress changes that occur for the scope of this work, but they do not capture the exact locations of the inflections and the maximum/minimum values, which can deviate over 100\,MPa from those originally measured. 


\section{Conclusions}
A study of the residual stresses produced following the welding of a series of ferritic steel plates with low transformation-temperature filler alloys has been performed using neutron diffraction. From this study the following conclusions and recommendations have been drawn:

\begin{itemize}

\item There appears to be no additional benefit to stress relief by using a highly enriched welding filler (LTT-2) to compensate for compositional dilution with the base plate for multi-pass welds. This is due to an overly suppressed $M_\mathrm{S}$ such that the $\gamma \to \alpha'$ transformation cannot be fully exploited at ambient temperatures in the final passes.

\item The use of a highly enriched welding filler (LTT-2) as a final capping pass successfully imparts minor compression into the surface regions of the deposited weld metal, however, large regions of high tensile stresses remain in the surrounding HAZ.

\item Compressive stresses were generated in the weld metal for a series of base plates that have differing yield strengths, above and below that of the LTT filler. However, residual stresses approximately equal to the yield strength of the base plate remained in the regions surrounding the HAZ. The re-austenised base plate material surrounding the fusion boundary will inevitably transform to martensite at an elevated temperature, which will lead to the accumulation of tensile stresses due to the high $M_\mathrm{S}$ of the plate.

\item The present work indicates that higher spatial-resolution ($\approx 1$\,mm) measurements of the stresses are necessary in order to adequately  capture the positions of inflection and the maximum/minimum stresses across the fusion boundary and HAZ.

\item It is believed that weld 1 provides the optimum stress distribution based on the maximum tensile stresses generated with respect to the yield strength of the base plate and the relative sizes and positions of these high tensile regions. Although weld 2 produces a comparable stress distribution, the retained austenite in the final passes will reduce the strength of the weld metal in this region. However, this maybe beneficial if a particularly tough and ductile weld is preferred.

%\item If more uniform mechanical properties are desired throughout the depth of the weld deposit, the appropriate procedure may be to begin welding with the highly enriched (LTT-2) alloy followed by depositing a welding filler with a lower alloying content (LTT-1). This deposition procedure would also have the effect of permitting the majority of the austenite to transform into martensite at the point where the weld metal reaches ambient temperature and theoretically maximising the stress-relief benefits of transformation plasticity.

\end{itemize}

 

\section*{Acknowledgments}
{\footnotesize We are grateful to ESAB AB for sponsoring this research. This work is based upon experiments performed on the ENGIN-X instrument at ISIS, Rutherford Appleton Laboratory, Didcot, U.K. (RB 1220125) and the STRESS-SPEC instrument operated by FRM II at the Heinz Maier-Leibnitz Zentrum (MLZ), Garching, Germany.}

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\end{thebibliography}

 
 
\clearpage

 

\begin{table}[!htbp]
\centering
\begin{center}
\caption{Combinations of welding consumables and base plates.}
\begin{tabular}{clccc}\\
\toprule \\
    Weld & Base Plate & Pass 1 & Pass 2 & Pass 3 \\[10pt]\midrule
      1          & Weldox 700 & LTT-1& LTT-1& LTT-1                          \\
     2         & Weldox 700 & LTT-2& LTT-2& LTT-2                             \\
     3         & Weldox 700 & Coreweld 89 & Coreweld 89  &LTT-2               \\
     4        & Weldox 700 & Coreweld 89& Coreweld 89& Coreweld 89                       \\
     5         & Weldox 960 & LTT-1& LTT-1 & LTT-1                             \\
     6         & 355 & LTT-1& LTT-1& LTT-1                       \\[10pt]
\bottomrule 
\end{tabular}
\label{tab:samples}
\end{center}
\end{table}


\begin{table}[!htbp]
\caption{Compositions (wt\,\%) and $M_\mathrm{S}$ of the undiluted welding alloys and base plates. The LTT  $M_\mathrm{S}$ were calculated based on a measured composition.}
\centering
\begin{center}
\resizebox{0.9\textwidth}{!}{\begin{tabular}{lccccccc}
\toprule 
Material & C& Si& Mn&Cr&Ni&Mo & Calculated $M_\mathrm{S}$ ($^\circ$C)\\
\midrule 
LTT-1 & 0.01-0.03 & 0.6-0.8 &1.2-1.7 &12.5-13.0 & 5.5-6.5 & $<$0.1 & 169\\
LTT-2 & $<$0.02 & $<$1 & $<$2 & 15-18 & 6-8 & $<$0.1 & 87 \\
Coreweld 89 & 0.12 & 0.65 & 1.50 & 0.70 & 2.80 & 0.85 & 372 \\
Weldox 700 &0.15&0.29&0.98&0.25&0.043&0.15 & 447\\
Weldox 960 & 0.17 & 0.22 & 1.24 & 0.2 & 0.052 & 0.65 & 420\\
355 & 0.16 & 0.26 & 0.95 & 0.012 & 0.017 & $<$0.01 & 451\\
\bottomrule 
\end{tabular}}
\label{tab:compositions}
\end{center}
\end{table}


\begin{table}[!htbp]
\centering
\begin{center}
\caption{Mechanical properties of the undiluted welding alloys and base plates.}
\resizebox{0.9\textwidth}{!}{\begin{tabular}{lcccc}
\toprule 
 Material & 0.2\% stress & Tensile strength & Elongation & Impact energy\\
   & (MPa) & (MPa) & (\%) & (at -20$^\circ$C) \\
\midrule 
LTT-1 & 627 & 1111 & 14.0 & 45\\
LTT-2 & 316 & 845 & 31.8 & 88\\
Coreweld 89 & $>$900 & $>900$ & 19.0 & $>$90\\
Weldox 700 & $>$700 & 780-930 & 14.0 & $>$27\\
Weldox 960 & $>$960 & 980-1150 & 12.0 & $>$27\\
355 & $>$355 & 470-630 & 22.0 & $>$27\\
\bottomrule 
\end{tabular}}
\label{tab:mech_prop}
\end{center}
\end{table}


\begin{table}[!htbp]
\centering
\begin{center}
\caption{Experimental parameters for stress measurements by neutron diffraction.}
\begin{tabular}{lcc}
\toprule 
Parameter & ENGIN--X & STRESS--SPEC\\
\midrule 
$\lambda$ & Time-of-flight & $\sim$1.69\AA\\
Diffraction peak & Multiple & \{211\} \\
2$\theta_\mathrm{M}$ & 76-104$^\circ$ & 77.3$^\circ$\\
Longitudinal $V_v$ & $3\times 3\times 3\,\rm{mm}^3$ & $2\times 2\times 2\,\rm{mm}^3$\\
Transverse $V_v$ & $3\times 15\times 3\,\rm{mm}^3$ & $2\times 20\times 2\,\rm{mm}^3$\\
Normal $V_v$ & $3\times 15\times 3\,\rm{mm}^3$ & $2\times 20\times 2\,\rm{mm}^3$\\
\bottomrule 
\end{tabular}
\label{tab:expt_parameters}
\end{center}
\end{table}


\begin{figure}[htdp]
\begin{center}
\includegraphics[width=0.5\textwidth]{plate.eps}
\caption{Machined plate set-up for welding.}
\label{fig:plate}
\end{center}
\end{figure} 


\begin{figure}[htdp]
\centering
\subfigure(a){
\includegraphics[width=0.6\textwidth]{comb_pos.eps}}\\
\subfigure(b){
\includegraphics[width=0.6\textwidth]{comb.eps}}
\caption{(a) Plan view of a welded plate showing the locations of the residual strain measurements and reference strain-free specimens. (b) Sectioned specimen used to determine the strain-free lattice parameter \cite{Ramjaun:2014}.}
\label{fig:weldcomb}
\end{figure}



\begin{figure}[!ht]
\begin{center}
\includegraphics[width=1.2\textwidth]{contour.eps}
\caption{Residual stress contours, in MPa, for the six different welds.}
\label{fig:contour}
\end{center}
\end{figure} 


\begin{figure}[!ht]
\begin{center}
\includegraphics[width=0.6\textwidth]{stress_plots.eps}
\caption{Longitudinal residual stress profiles as a function of distance from the weld centreline.}
    \label{fig:stress_plots}
    \end{center}
\end{figure} 


\begin{figure}[!ht]
\begin{center}
\includegraphics[width=0.8\textwidth]{stress_temp.eps}
\caption{Schematic illustration of the residual stresses for three welding alloys with differing $M_\mathrm{S}$, highlighting the effect of an overly suppressed $M_\mathrm{S}$.}
    \label{fig:stress_temp}
\end{center}
\end{figure}


\begin{figure}[!ht]
\begin{center}
\includegraphics[width=0.8\textwidth]{Fmacro.eps}
\caption{Macrostructure of weld 2 revealed using ferric chloride.}
    \label{fig:Fmacro}
\end{center}
\end{figure} 


\begin{figure}[!ht]
\begin{center}
\includegraphics[width=0.8\textwidth]{CFGH_profile.eps}
\caption{Welds 1, 2, 3, 4: Averaged longitudinal residual stress profiles at 2.5\,mm below the top surface as a function of distance from the weld centreline.}
    \label{fig:CFGH_profile}
\end{center}
\end{figure}


\begin{figure}[!ht]
\begin{center}
\includegraphics[width=0.8\textwidth]{extra.eps}
\caption{Weld 5: Longitudinal, transverse and normal residual stress profiles as a function of distance from the weld centreline highlighting the change in profile with additional strain measurement positions at 5\,mm below the surface.}
    \label{fig:extra}
\end{center}
\end{figure}


\end{document}
