**1. Introduction**

It has been over 61 years since the first production heat of a niobium (Nb) bearing HSLA strip steel was commercially produced in 1958 [1,2]. Since that time, microalloying with Nb has been extended to virtually all product classes and forms. Perhaps nowhere has Nb been more beneficial than in linepipe steels for high pressure oil and gas transmission at low temperatures. The contribution of Nb to high strength by grain refinement; solid solution strengthening of the various types of ferrite formed; and further strengthening of the ferrite by precipitation and dislocation-hardening through increased hardenability have all been chronicled in numerous papers and research studies [3–8]. In a similar fashion, the improvement in lowering the ductile-brittle transition temperature (DBTT) through the use of niobium is critically important in pipelines intended for low temperature service [9]. Much of this improvement is caused by the grain refinement of both the austenite during hot rolling, and the grain refinement of the ferrite during transformation upon cooling. Several studies have shown that strength and toughness can be improved by increasing the level of Nb used, from zero to the

conventional 0.03–0.04 wt.%, to very high levels near 0.1% Nb [10,11]. However, there are detriments beyond cost to the use of Nb, especially at higher levels, such as reduced weldability, large undissolved particles and higher mill loads.

When a 25 mm thick plate intended for pipe applications is hot rolled on a modern 5 m plate mill, it does so in three stages after slab reheating: rough rolling, finish rolling and accelerated cooling. For example, after reheating at 1200 ◦C, a continuously cast slab of 200 mm thickness might be rolled to 75 mm transfer bar in the roughing mill at temperature from 1150–1050 ◦C. The bar is then transported to the finishing stands, where it is further rolled to the final gauge of 25 mm at temperatures ranging from below 900 to 750 ◦C. Because of the large width and potentially high mill loads, the rough rolling and finishing rolling take place in numerous, light passes, often between 10 and 15 passes in each case, with pass reductions between 10–15% each in the roughing mill, and 5–10% each in the finishing mill. Since these are reversing mills, the interpass times can range from 10–30 s depending on conditions.

After final rolling, the plate is water spray cooled at about 30 ◦C/s to the water end temperature (WET) which is dictated by the continuous cooling transformation (CCT) diagram, the cooling path and the required microstructure and properties [9]. The WET for achieving ferrite–pearlite microstructures in steels of high hardenability, i.e., CEII over 0.40, is in the range of 600–650 ◦C. When a bainitic ferrite microstructure is desired, a WET near 450–550 ◦C would be used. For even higher strengths, requiring martensite, a WET around 300 ◦C would be used. Normally, the steels are air cooled to room temperature from the WET.

The microstructure of austenite during hot rolling of Nb-bearing HSLA steels is believed to behave according to Figure 1 [12]. For a given strain, above *T*95%, austenite grains are repeatedly recrystallized and grain refined (complete recrystallization; RXN), while below *T*5%, the austenite grains are unrecrystallized and pancaked. These critical temperatures are themselves dependent on the content of Nb of steels and increase with higher Nb levels, as shown schematically in the red-dashed curves in Figure 1.

**Figure 1.** Schematic diagram of the resulting austenite microstructures from different deformation conditions, after reference [12].*T*5% and*T*95% are temperatures for 5% and 95% recrystallization, respectively.

The hot processing of Nb steels is based on most of the Nb being in solution in the austenite during slab reheating and rough rolling, and some of the dissolved Nb being reprecipitated as strain-induced precipitate in deformed austenite during finish rolling. The nose of the C-curve for this precipitation in austenite is around 900 ◦C for 0.04 wt.% Nb, but is raised to 950 ◦C or higher at larger levels of Nb [3,13]. As mentioned above, studies of the recrystallization-stop-temperature, *T*5% or *T*nr, indicate that this temperature increases with Nb content and can reach or exceed 1050 ◦C in typical pipe steels containing 0.1% Nb [6,14]. Since the last roughing passes occur in this temperature range, it is quite possible that there might be strain-induced precipitation of NbC, even in the late roughing passes. Since the grain refinement expected during roughing requires multiple static recrystallization events in the interpass times, perhaps complete recrystallization may not occur by the exit of the roughing mill. Therefore, there might be consequences for the grain refinement needed during rough rolling when the bulk Nb level is too high. For example, this lack of sufficient grain refinement is important, since coarse grained microstructures often result in failure in drop weight tear testing of plate and pipe.

Whether static recrystallization occurs or not depends on the comparison of the driving force for static recrystallization and the retarding Zener pinning force caused by the interaction of the moving austenite boundaries with particles formed earlier on the defect structure of the deformed austenite in the roll gap; these defects include grain and subgrain boundaries and deformation bands [15].

The potential for strain-induced precipitation of NbC or NbCN will depend on the bulk composition of the steel, the relevant solubility products for NbC or NbCN in austenite and the applicable rolling practice. The carbon content will be governed by the final strength needed in the final pipe, which can be X50–60 for ferrite–pearlite (F–P) microstructures, X70–100 for bainitic microstructures and X120 and above for martensitic microstructures [16,17]. While changes in %C have only a slight influence on the strength of F–P steels, they can have dramatic effects on the strength of bainite or martensite found in direct quenched steels. Therefore, the carbon content can be expected to influence both the strength of the plate or pipe through hardenability effects, and the toughness through its effect on the precipitation in austenite and the control of grain size.

The effect of hot rolling on toughness or lowering the DBTT, through grain refinement of the prior austenite grain size (PAGS), is the result of two sequential events. The first is the combination of the elimination of both the remaining as-cast structure and the large grain size that result from slab reheating. These occur during the repeated static recrystallization that takes place between the rough rolling passes that arise when the pass strains happen above the *T*95% temperature. This leads grain refinement during multiple waves of recrystallization, where the PAGS might be reduced from 300 μm to 50 μm going into the finishing passes where pancaking occurs.

Since the finishing passes for controlled rolling are usually considered to occur below the *T*5% or recrystallization stop temperature, normally about 900 ◦C in a steel containing 0.04% Nb, the final austenite will be heavily strained, elongated or pancaked [18]. This austenite is highly deformed with high strength and contains numerous crystallographic defects, such as deformation twins, deformation bands, subgrain boundaries and elongated grain boundaries. These near-planar crystalline defects contribute to what is call the *S*v value, an index used to judge the effectiveness of a given rolling process in thermomechanical processing (TMP) [19]. In low hardenability or slowly cooled F–P steels, the nucleation of polygonal ferrite occurs on the *S*v, where the high density of high angle ferrite grain boundaries resulting from ferrite grain refinement can act as crack arresters for potentially growing cleavage cracks. In higher hardenability or faster cooled steels, the defects themselves can act as cleavage crack arresters in bainitic or martensitic steels. In either case, these near planar defects can act as sites for crack arresters for the growth of cleavage cracks; hence, lowering the DBTT.

To investigate the rough rolling process, the possibility of strain-induced precipitate was explored in this research in the temperature range for rough rolling, i.e., 1150–1000 ◦C, for a series of steels with two carbon levels: a conventional one of 0.06 wt.% C and a lower carbon version of 0.03 wt.% C. Three levels of Nb were also studied—0, 0.04 and 0.08 wt.%.

The purpose of the current study was threefold: (i) to determine whether NbC could form during rough rolling and (ii) whether this precipitate might cause less than complete recrystallization during rough rolling, and (iii) therefore, lead to a larger slightly pancaked as-roughed austenite grain size resulting in non-optimum final austenite microstructure prior to transformation. To achieve these goals, two investigations were conducted. The first was a theoretical study of potential strain induced precipitation of NbC in the temperature range in which roughing passes would normally occur in a modern 5 m wide reversing plate mill. This required solubility calculations, the use of the subgrain boundary Zener pinning model and parameters taken from similar, earlier studies [3,18,20,21]. The second was the study of the austenite grain size and shape in quenched specimens after laboratory hot deformation experiments conducted in the temperature range in which rough rolling would normally occur. It was expected that this research would help answer the question as to whether high Nb steels can be expected to be successfully processed on modern plate mills in a range of strength levels.

#### **2. Research Rationale**

## *2.1. Alloy Design Used in Study*

The experimental alloys were chosen to represent a typical linepipe steel, near X100 composition [16]. As mentioned earlier, the carbon and niobium levels were chosen to demonstrate the effects of NbC pinning forces in the temperature range of rough rolling. Figure 2 shows the temperature profile of a modern, 5 m wide reversing plate mill where the end of roughing rolling is approximately at 1050 ◦C [22].

**Figure 2.** Temperature profile of a modern 5 m wide reversing plate mill, after reference [22].

NbC in solution was calculated using the equilibrium solubility product proposed by Palmiere et al. [20]:

$$\log[\text{Nb}][\text{C}^{\cdot}] = 2.06 - \frac{6700}{\text{T}} \tag{1}$$

Figure 3 shows the calculated dissolution temperatures of alloys with two Nb levels, namely, 0.04Nb and 0.08Nb, in wt.% while keeping carbon as a constant at 0.06 wt.%. It is clear that all the NbC available in 0.04 wt.% Nb steel will be dissolved at/or close to 1170 ◦C, while the higher Nb content will not be fully dissolved below a slab reheating temperature of 1265 ◦C, which indicates more stable precipitates at higher temperatures. It should be noted that this temperature far exceeds the normal reheating temperature of 1150–1200 ◦C used in practice. The dashed line represents the deviation between the two alloys in terms of soluble NbC; thus, the reheating temperature is not a variable, but is restricted to 1150–1200 ◦C in practice. Therefore, the question is how much Nb can be dissolved at perhaps 1200 ◦C. According to Figure 3, all of Nb in the 0.04 steel and 0.55% Nb in the 0.08 steel would be dissolved at 1200 ◦C. In normal practice with a commercial steel of similar composition containing 0.04Nb, all of the Nb is taken into solution during reheating, and very little if any is reprecipitated during rough rolling. Therefore, recrystallization goes to completion in roughing. However, in the 0.08Nb steel, only 0.055 wt.%. Nb is taken into solution and the remaining 0.025 wt.% Nb is left undissolved. The question is: will any of the 0.055 wt.% Nb dissolved in austenite reprecipitate during roughing, and if it does, will it suppress complete recrystallization during roughing? Higher Nb levels

indicate that some precipitates could not be taken into solution based on solubility considerations in the range applicable to the rough rolling process.

**Figure 3.** Calculated amount of dissolved Nb at equilibrium as a function of reheating temperature.
