physical understanding of ferrite grain growth during welding ...

E. Bayraktar, J.P. Chevalier, D. Kaplan, and L. Devillers

June 2009 The Arabian Journal for Science and Engineering, Volume 34, Number 1C 115

PHYSICAL UNDERSTANDING OF FERRITE GRAIN GROWTH DURING WELDING IN INTERSTITIAL FREE

STEELS (IFS) +

E. Bayraktar*

1Supmeca-Paris, School of Mechanical and Manufacturing Engineering, France

J. P. Chevalier 2Chair of Industrial Materials, CNAM, case 321, 292 rue Saint Martin, 75141 Paris 03 –

France

D. Kaplan and L. Devillers

ARCELOR Group, Paris-France

:ةـالخالص

في دراستنا ) IFS(خالل تلحيم مواّد الفوالذ الخالي من الفجوات ) HAZ(تّم التقرير عن الُنمّو الُحَبيبّي غير الطبيعي في المنطقة الُمتأثرة بالحرارة وتحصل هذه . جاهات انسياب الحرارةوقد أظهرت مالحظة الُوصالت الُملّحمة وجود ُحبيبات آبيرة جدًا قرب خط االنصهار وُموّجهًة مع ات]. 1[السابقة

Ac3 معّينة على مسافة ُمعّينة من خط االنصهار وتتوافق مع المناطق التي تزيد درجة الحرارة الُقصوى الحاصلة فيها قليًال على IFSالظاهرة في مواّد

150µmالُنمّو الُحَبيبّي ولدرجات مختلفة من الفوالذ بين وتراوحت األحجام الُمقاسة لُحبيبات الِفّريت في منطقة . γ α)درجة حرارة التحّول من (ُمعّدل التدّرج الحراري الذي : وتّم تعليل هذا الُنمّو الُحَبيبّي الغير طبيعي باستعمال عاملين هما]. 2، 1[ وذلك بحسب أحوال عملّية اللحام 500µmوحتى

ومن الواضح أّن النمّو . من خالل المنافسة بين التنّوي والُنمو في تدرج حراري) Ar3 )R ُيشّكل الُقّوة الدافعة، ونسبة زيحان خط تساوي درجة الحرارة . غير ُمعّوق بوساطة الشوائب في المواّد الُمترّسبة γ αالسريع لُحبيبات الِفّريت يمكن حصوله إذا آان التحّول

.لُنمّو الُحَبيبّي غير الطبيعي وسط تدرج درجة الحرارة الحاصل خالل عملية اللحامُتقّدم هذه الورقة البحثية أول وسيلة توضح أثر حالة الترّسب على ا

ABSTRACT

Abnormal grain-growth in the heat affected zone (HAZ) during the welding of interstitial free (IFS) steels has been reported in our previous study [1]. Observations of the welded joints indicate the presence of very large grains near the fusion line, and these are oriented along the directions of heat flow. This phenomenon takes place in certain IF steels at a given distance from the fusion line and corresponds to regions where the maximum attained temperature is slightly over Ac3 (γ α transformation temperature). In the grain-growth zone, the measured ferrite grain size for different steel grades varies from 150µm up to 500µm according to the welding conditions [1,2]. This abnormal grain-growth was interpreted by using two parameters, the thermal gradient (G), which constitutes the “driving force”, and the displacement rate of the Ar3 isotherm (R) in terms of competition between nucleation and growth in a thermal gradient. Clearly, the rapid growth of ferrite grains can occur only if the γ α transformation is not impeded by impurities in the precipitates.

This paper presents a first approach demonstrating the influence of the precipitation state on abnormal grain growth in temperature gradient produced during welding.

Key words: welding of IFS, ferrite grain growth, HAZ, precipitation effect

∗Corresponding Author: E-mail: [email protected]

This paper was presented at the Advanced Manufacturing Processes and Technologies (AMPT) Conference held in Bahrain on November 2–5, 2008.


The Arabian Journal for Science and Engineering, Volume 34, Number 1C June 2009 116

PHYSICAL UNDERSTANDING OF FERRITE GRAIN GROWTH DURING WELDING IN INTERSTITIAL FREE STEELS (IFS) +

1. INTRODUCTION

Ultra-low carbon interstitial-free (ULC, IFS, <15-20 ppm) steels are manufactured by adding carbide, nitride, and sulphide-forming elements, such as titanium, boron, and or niobium, in a “RH-OB” vacuum degasser. In the case of Ti added IF steels, the carbon, nitrogen, and sulphur elements form precipitates of different types, such as TiN, TiS, Ti4C2S2, and TiC. The other alloying elements (Mn, P) can also form precipitates of MnS and Ti–Fe–P types. The volume fraction of these precipitates is observed to be extremely low (~ 10-4 - 10-5). Additionally, these precipitates are relatively coarse (10 – 400 nm). The increase in precipitation hardening due to these precipitates is of the order of 1–5 MPa and is not worth mentioning in IF steels [1–11].

In the literature, a common idea is that all of these precipitates cause a complete elimination of nitrogen, carbon, and sulphur from the solid solution. Thus, one possible way to improve texture is to decrease interstitial elements as much as possible in solutions for steels suitable for very high quality deep drawing properties (hyper deep drawable steels in the car industry). Each processing step influences the microstructural evolution of cold rolled and annealed sheets of high formability. IF steels develop a very strong orientation of texture with (111) <110> parallel to the sheet surface, resulting in a high anisotropy (Lankford) ratio « r ≥ 2 » and very good deep drawing behavior. For this, four principal factors are assumed to play an important role; these are the absence of carbon in solid solution, the influence of the different elements, the grain size, and the influence of the precipitates, their size, and particularly their quantity [11–14]. Thus, it is important to control the precipitate characteristics (type, amount, size, and distribution), as they strongly influence the properties of IFS, and it is also necessary to diminish the quantity of interstitial elements in the solution matrix, mostly carbon and nitrogen, in order to obtain exceptional ductility and deep drawing proprieties. The interstitial elements in IF steels are stabilized by means of the addition of some micro alloying elements, such as titanium and niobium, which can trap nitrogen, sulphur, and carbon very easily in the stable precipitated forms during hot rolling. Due to the weak percentages of C and N obtained, the percentages of Ti and/or Nb required are, therefore, very low for these grades of steels. Additionally, these added (stabilizing) elements in IF steels decrease the recrystallization rate and annealing temperatures. Due to the total trapping of C in the form of TiC and/or NbC in IF steels, they are insensitive to ageing and, thus, can be produced directly on rapid annealing lines that do not dispose of the over aged section. It is also unnecessary to perform a post-annealing of the over aged section in compact coils. However, an incomplete trapping of C in the IF-TiNb grades can bring back the bake hardening properties without damage of the "r" value. Moreover, the annealing at recrystallization aims to give to the steel an optimal recrystallized grain size for obtaining a reasonable level of yield strength, and, for the deep drawing quality steels, to develop favorable crystallographic orientations for obtaining a high anisotropy ratio « r ». Use of very low carbon IF steels stabilized with Ti and/or Nb leads to hyper-deep drawable steels produced continuous annealing with significantly higher mechanical properties [1–5,9,11–15].

In fact, precipitates can present resistance to dislocation motion by means of the Friedel effect (shearing of precipitates by dislocations) or the Orowan effect (evading of precipitates by dislocations) mechanism. This resistance leads to a considerable increase in strength and is strongly governed by the size, spatial distribution and volume fraction of the precipitates [14–17]. It should be noted that IF steels contain very small amounts of alloying elements.

The grades of IF steels are classified in different categories: IF-Ti, IF-TiNb and/or IF-TiNb, IF-TiNb-B, IF-B, and IF-Ti-B. In fact, titanium in IF-Ti grades is added in quantities larger than that required for the precipitation of stochiometric TiC [1]. An addition of niobium also traps the available carbon [1–5]. The co-addition of Nb with Ti can very easily decrease the recrystallization temperature. Moreover, an addition of boron is an effective method of preventing abnormal grain growth in IFS during welding operations because B and B precipitation can be found in the intergranular ferrite plates in a coarse grain [1,16–22]. This phenomenon takes place in certain IF steels at a given distance from the fusion line and is related to regions where the maximum temperature attained is slightly over the α → γ transformation temperature, Ac3 [1,12,14,21]. Moreover, the phenomenon of abnormal grain growth was primarily interpreted in relation to the thermal gradient, which must be the “driving force” [1]. However, it can be assumed that the abnormal grain growth can occur only if its development is completed easily within a matrix of great purity and/or very low precipitates. Keeping in mind this claim, one observes indeed that the phenomenon of abnormal grain growth does not occur during the welding of steel grades, in which the carbon content is slightly higher than that of IFS (example: C > 15-20.10-3 %), and under the same conditions of a very intense thermal gradient (resistance spot welding - RSW, GTAW, LASER welding, etc.). Thus, this stage of the research requires an understanding of the role and/or the effect of each precipitate individually on the ferrite grain growth phenomenon during the welding of the various grades of IFS examined here.



2. EXPERIMENTAL CONDITIONS

Different industrial grades of IF steels were studied. Table 1 presents the compositions of these steels, with thickness varying between 0.7 and 0.9 mm. Grain growth in the HAZ was studied using the following processing conditions:

- Gas tungsten arc welding (GTAW) was used with a voltage of 10V, an intensity of 65A, a welding rate of 40cm/min, and argon gas shielding (12 l/min). Considering the coefficient of thermal efficiency of the process (η≈ 50%), the corresponding linear energy varies from 0.75 to 0.95kJ/cm.

- Resistance spot welding experiments were carried out with a voltage of 1V, an intensity of 6000A, and a welding time of 30 periods (50 Hz). The corresponding energy varies from 1.5 to 8.5 kJ. The spot diameter was approximately 4 mm.

- Grain size was measured using optical metallographic microscopy.

- TEM (transmission electron microscopy) observations were carried out on the base metal and welded specimens. TEM was carried out on 3 mm diameter foils. A Tenupol 6 twin jet electro polisher was employed with an electrolyte of 5 vol. % perchloric acid and 95 vol. % methanol. TEM examination of thin foils was performed using a JEM-2000-CX1 TEM microscope operated at 200 kV.

- Because of the considerable anisotropy of the ferrite grain sizes, the grains were measured individually in the longitudinal, dαL, and transverse, dαT, directions of the grain in order to derive a mean value in each direction.

- The calculations of precipitates in the different grades of IFS were carried out by using a special PC program developed by IRSID (ARCELOR).

Table 1. Chemical Compositions of Different Grades of IF Steels-Non Commercial (10-3 wt %)

Steel grade C Mn P S Si Al Ni N Cr Cu Nb V Ti B Ceq IF Ti) (0.8 mm)

1.4 205 13 10 4 38 33 3.5 20 13 2 3 107 44

IF Ti (0.9 mm)

5 195 6 9 7 32 20 3.8 15 9 1 2 90 44

IF-TiNb (0.8mm)

3.3 140 13 6.6 20 7 21 20 7 18 18 36

IF TiNbB (0.7mm)

3.1 158 7 5.9 11 49 17 3.4 19 10 18 26 0.6 37

IF TiB (0.67 mm)

2.7 159 8 7 5 37 17 2.4 16 4 75 10 35

3. INITIAL STATE OF PRECIPITATION (BEFORE WELDING) IN INTERSTITIAL FREE STEELS (IFS)

Many direct observations have been carried out for the IFS grades [1–14, 16–21]. However, many of these reports used hot deformation (rolling) conditions, and only a few studies were published relating to real welding operations. In this paper, for the sake of simplicity, the state of precipitation was first estimated by means of calculations of the products of solubility. After that, the evolution of the state of the precipitate fraction volume in HAZ was determined for each grade as a function of the temperature reached during the welding operation. In a general way, the molar fraction of the precipitated phase in equilibrium can be estimated by the product of solubility:

3.1. Log (%M) (%C, N) = A - B/T

where M and C are the percentages (wt % or at %) of the elements liable to form the compound,

A and B are coefficients used in this calculation and T is the solution temperature (in Kelvin).

T is the temperature of dissolution (in Kelvin).

By detailing the various grades of IFS studied here, the precipitates were taken into account separately for each grade, such as IF-Ti, IF-TiNb, and IF-TiB in this analysis.

Thus, the objective of this study was to research the factors (steel grades and welding conditions) and clarify the role of each precipitate on the abnormal grain growth phenomenon qualitatively. A typical study that is of special interest in the manufacturing of a body of a car was then undertaken on TIG (GTAW) and resistance spot welding (RSW) of the different grades of IFS, and a special application of tailored blanks was given on the grade of IFTiNb.



The analysis of the precipitates in the various grades of IFS studied are summarized as follows:

3.2. In the Case of IF-Ti Grade

One primarily finds in this type of grade:

- Titanium nitrides (TiN)

- Titanium sulphides (TiS), of which some are most probably transformed into carbon-disulphides Ti4C2S2 at the coiling stage in the presence of free carbon.

- Titanium carbides (TiC) of 10–20 nm (Intragranular) or 50 nm (Intergranular) precipitating during hot rolling and mainly at the coiling stage.

3.3. In the Case of IF-TiNb Grade

The precipitates of the Nb in the form of unstable sulphides of the considered contents form, and only the type of compounds known as carbon-nitrides, facilitated by possible substitution of nitrogen and carbon, titanium, and niobium. These precipitates can form in these grades during hot rolling and coiling. Nitrogen can be stabilized with Ti. These compounds can, thus, be described in the general form: (TiαNbβ) (NXC1-x). The following calculations, thus, consider the formation of (Ti, Nb) N (size: 40-100 nm), TiS (size: 500 nm), and NbC (size: 5 -40nm). They show a higher density of small precipitates (108 ppt/mm2) compared to IF-Ti steel (106 ppt/mm2).

3.4. In the Case of IF-TiNbB Grade

One can consider that the precipitates to be taken into account are the same ones as those in the steel grades of IF-TiNb, in so far as boron cannot form any precipitates with the nitrogen or the carbon, which are already fixed by other elements, such as titanium.

For the sake of simplicity, calculations have been carried out by supposing a co-precipitation of the elements, because the different precipitates are not soluble in each other. Table 2 specifies the values of the expressions of the products of solubility in the austenite, as well as the values of the enthalpy of the different compounds at ambient temperature.

For the compositions studied, Table 3 shows the quantities of the different elements precipitated at ambient temperature. It should be noted that titanium is present in over – stochiometric values1 compared to the nitrogen in the grades of IF-Ti and IF-TiB. The nitrogen is, thus, entirely precipitated in the form of TiN. The excess of titanium also can fix the carbon in these grades. This case is valid only in the grades containing low titanium content (IF-TiNb, IF-TiNbB, etc) so that only a part of carbon is precipitated in the form of NbC.

4. EVOLUTION OF THE PRECIPITATES IN HAZ DURING WELDING

Following the example given in [1], an assumption will be made here that the evolution of the initial precipitates can be evaluated by means of the products of solubility. This step does not take into account the kinetic problems naturally; the very high heating rates in welding delay the dissolution towards higher temperatures. This type of approach, however, makes it possible to obtain default values of the real temperatures of dissolution.

Table 4 shows the calculated temperatures of complete dissolution of the precipitates as discussed in the previous section. According to grades studied here, these temperatures are around 1360–1520°C (TiN), 1140–1300°C (TiS), 800–960°C (TiC), and 570–720°C (NbC). By order of increasing stability, one finds order of classical development: carbides< sulphurs< nitrides.

Table 2a). Coefficient Values of the Products of Solubility in Austenite Used for Calculations of the Precipitates, b) Free Enthalpy Values of the Different Simple Compounds at Ambient Temperature

Precipitates A B Field of calculationTiN 5.19 15490 austenite TiS 8.2 17640 austenite Tic 5.02 10500 austenite

NbC 2.81 7019 Ferrite

a)

1 The excess of titanium in solution allows for a very favorable increase in the deep drawing capacity of steel

by delaying of the restoring phenomenon of work hardened structures.



∆G kJ/mole

∆H kJ/mole

Carbide Sulphur Nitride

-149 to 168 -160 -200 -235

-134 to -152-201 -338 -297

NbC TiC

TiS

TiN

b)

Table 3. Percentage (10-3 wt %) of Precipitate of Each Element at Ambient Temperature in the IFS as Received Conditions

Moderate Precipitate Ti Nb S C N IF-Ti

(0.8mm) TiN TiS TiC

11.9715 5.6

- 10

1.4

3.5

IF-Ti (0.9mm)

TiN TiS TiC

13 13.4 20

9

5

3.8

IF-TiNb (0.8mm)

TiN TiS NbC

10.267.74

18

5.18

2.32

3

IF-TiNbB (0.7mm)

TiN TiS TiC NbC

11.638.81 5.52

13.32

5.9

1.381.72

3.4

IF-TiB (0.67mm)

TiN TiS TiC

8.9 10.5 10.4

7

2.6

2.6

Table 4. Temperature of Dissolution to Balance of the Various Compounds Formed in the IF Steels Studied

PrecipitateIF-Ti

(0.7mm) °C

IF-Ti (0.8mm)

°C

IF-Ti1289(0.9mm)

°C

IF-TiNb(0.8mm)

°C

IF-TiNbB(0.7mm)

°C

IF-TiB (0.67mm)

°C TiN 1428 1525 1516 1365 1403 1466 TiS 1274 1306 1274 1139 1164 1264 Tic 890 914 960 865 800 935

NbC - 566 592 725 669 -

Moreover, the progressive dissolution in solid solution of the various compounds was evaluated by calculating, from the various products of solubility and of the chemical compositions, the precipitated fraction depending on the temperature reached in the HAZ. The results are shown in Figure 1, which presents examples of such a calculation for the grades IF-Ti and IF-TiNb.

Figure 2 also gives the evolution of the stability of different compounds in IF-TiNb and the estimation of the solid solution as a function of the temperature. In other words, it gives the details of the evolution of the stability of different components and estimations of probable dissolution as a function of temperature.

In the case of this steel (C= 5.10-3 %, N=3.8, S= 9, Ti=90), this type of presentation verifies the different stages of the dissolution: initially, the titanium carbides, the most present precipitates, are partially decomposed before the γ α transformation and completely disappear even in the HAZ at low temperature (950–1000°C). This is only valid for the higher temperatures (1100°C) where the titanium sulphides themselves are also given in solid solution; in another words, this stage becomes complete around 1300°C. In the same way, TiN also begins to decompose at this temperature.

From now on, the interaction level of these various phenomena will be examined with the grain growth mechanism that was reported in reference [1].



400

500

600

700

800

900

1000

1100

1200

1300

1400

1500

1600

0 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08

IF-Ti (1289)wt %(10-3)

C : 5, S : 9, N : 3,8, Nb : 1, Ti :

TiN

TiS

Θ(°C)

TiC

precipitate frac, volume (%)

γ

α

a)

Figure 1a. Evolution of the state of the precipitated fraction in the HAZ of a grade of IF-Ti and b) IF-TiNb depending on the temperature reached during welding

400

500

600

700

800

900

1000

1100

1200

1300

1400

1500

1600

0 0.01 0.02 0.03 0.04 0.05 0.06 0.07 0.08

°C

TiN

precipitate frac, volume (%)

NbC

TiS

IF-TiNbwt %(10-3)

C : 3,3, S : 6,6, N : 3, Nb : 18, Ti : 18

Figure 2b. Evolution of the stability of different compounds in the grade of IF-TiB as a function of the temperature

5. INTERACTIONS BETWEEN THE PHENOMENON OF DISSOLUTION OF THE PRECIPITATES AND THE GROWTH OF GRAINS

There are many studies concerning the influence of the state of precipitation on the grain growth phenomenon in IFS during welding [6–12]. This interaction is typically described by a formalism more or less like that proposed by Zener, as given in the studies of [9–18], which expresses the fact that a population of many fine dispersed particles can break down the grain growth, since the displacement of the grain boundaries through these particles is accompanied by a local increase in free enthalpy. In fact, this description is not so significant and cannot be applied



in the case of the welding of IFS considered here, because it relates mostly to the interaction of the ferrite grain growth in an austenitic phase with the precipitates. The interpretation suggested here is based not on austenite grain growth but on the mobility of the interface γ α in the presence of a very huge thermal gradient [1]. The problem could, thus, be formulated as follows: what is the state of precipitation “met” during the displacement of the interface towards the weld metal bead?

As explained in the former work [1], the first ferrite germ was born in the joint welded towards 950°C in the cooling stage (at a certain distance from the fusion line). At this temperature, the niobium carbides are completely dissolved, and the precipitates of TiC are in partial dissolution. The displacement of the interface γ α towards the weld metal bead is illustrated schematically in Figure 3. Here the example was given for a steel grade of IF-TiNb. The interaction between the precipitations – the ferrite grain growth -- was schematically represented for the abnormal ferrite grain growth in the case of mono and two successive thermal cycles during the weld (successive positions, “A”, “B”, “C”).

In other words, the interface γ α moves gradually to the zones:

-Having undergone a cycle of austenitization at higher and higher temperatures

-Having undergone a cooling at lower and lower temperatures, the transformation point γ α is lowered with the maximum temperature.

On this last point, it is reasonable to imagine that the precipitation occurs much less during the cooling phase, taking into account the very fast cooling rates. Even in the case of a cooling rate of ∆t300

700 = 100s, many researchers have shown [9–15,18–23] that the fraction of precipitates was reduced during the cooling rate. In other words, the state of precipitation “met” by the transformation interface of α/γ is without doubt very close to that existing at the maximum temperature of the welding cycle. This means that the progression of the interface occurs within a structure containing fewer and fewer precipitates (the progressive disappearance of TiS, Ti4C2S2, TiN….) and this moves more and more towards the weld metal bead. Under this aspect, one will observe that the grade of IF-TiNb, which is most sensitive to the abnormal grain growth during the weld thermal cycle, presents a lower quantity of precipitates TiN and TiS and more precocious dissolution2.

Thus, in order to illustrate the possible role of the state of precipitation on the grain growth during the welding and to separate its role from that of the thermal gradient, some complementary tests have been carried out, which are described below.

6. INFLUENCE OF THE INITIAL STATE OF PRECIPITATION ON THE GRAIN GROWTH

6.1. Without the Presence of a Thermal Gradient

Certain cycles of reheating with a maximum temperature of 950°C have been carried out (temperature likely to reveal the ferrite grain growth) on a steel grade of IF-Ti (thickness = 0.7mm) in various metallurgical states:

- The first series of the test have been carried out as received conditions

- The second series of the test have been carried out after a preliminary heating to 950 or 1150°C3

These tests were carried out with a dilatometer, i.e. without a thermal gradient within test-tubes heated in a homogeneous way. The microstructures obtained are given in Figure 4. In the case of a thermal cycle starting from the received condition (initial state with precipitates of TiN, TiS, TiC), the grain size obtained is fine, on the order of 10 µm. It was seen previously that, in this temperature range, only the precipitates of the titanium carbides undergo a likely modification. The temperature 950°C corresponds to dissolution of these precipitates (total solid solution: 890°C). When the initial state of precipitation is modified (for example, with a preliminary cycle at 950°C), the grain size increased considerably (225 µm), which indicates that the titanium carbides play an important role in the grain growth. After a preliminary cycle at 1150°C (which, in addition to the titanium carbides, partially eliminates the precipitates of TiS, the temperature of complete dissolution of which is 1251°C), the grain size is 250 µm, which indicates that the contribution of these last precipitates to the resistance to the grain growth is relatively weak. As is well known, the precipitates of the TiN and TiS, formed at high temperature, are coarser than the others (size > 1µm) and they, therefore, have a very weak effect.

2 At 1100°C, for example, the TiS entirely disappeared in the grade of IF-TiNb, while it exist still in other grades 3 Thermal cycles are considered as the following: -One cycle: Vheating= 200°C/s, θM=950°C, dwell time=1s, Vcooling=500°C/s. -Double cycle: 1st one Vheating= 200°C/s, θM=950°C, dwell time=1s, Vcooling=500°C/s, 2nd one Vheating= 200°C/s,

θM=950/1150°C, dwell time=1s, Vcooling = 500°C/s.



TiS, Ti4C2S2

TiN

TiC, NbC

≈1300°C

≈1100°C

≈800-900°C

Θ(°C)

t(s)

Grain growth directionBM

A CB

950°C

1000°C 1200°C

Two successive thermal (weld) cycles

1st Cycle A 2nd Cycle

Θ (°C

t (s)

Figure 3. Interaction between the precipitations – the ferrite grain growth: Schematic representation of abnormal ferrite grain growth in the case of mono and two successive thermal cycles during the weld (example was given for a steel grade of IF-TiNb)

a) b) c)

Figure 4. Influence of initial state of precipitation on the grain size simulation in a dilatometer, steel IF-Ti, t=0.7mm: a) One thermal cycle θmax=950°C dα = 10µm, b) two successive thermal cycles θmax = 950+950°C, dα = 225µm and c) two successive

thermal cycles θmax =1150+950°C dα = 250µm



6.2. With Presence of a Thermal Gradient

In order to understand the influence of the initial state of the precipitation on abnormal grain growth in the presence of a large thermal gradient, the following experiments were carried out on several steel grades of IF-Ti and IF-TiNb: firstly, a GTAW remelting was carried out, which makes it possible to understand the grain growth obtained in the HAZ of welded steel sheets with an initial state of precipitates characteristic of industrial sheets. Then, a second GTAW remelting was carried out, so as to cross the first line of welding. This type of the test, thus, makes it possible to obtain a zone in which the thermal cycle associated with the HAZ of the second line of welding is applied to a state where the precipitates are largely decomposed by the first welding cycle. The grain size can be measured in this zone of intersection of the two weld beads and compared with the former grain size. Figure 5 indicates a typical microstructure observed in the zone of intersection (crossing zone).

All of the tests carried out in this study verify a very significant grain growth (for example, 400 µm in the case of a single remelting, with 800 µm in the zone of intersection), which correctly explains the importance of the initial state of precipitation.

In order to compare the results, all of the tests carried out on IF-Ti (t = 0.7mm) are summarized in Table 5, including the cases with and without a thermal gradient.

a)

b)

2nd fusion (Welding line)

1st fusion (Welding line)

2nd fusion (Welding line)

2nd fusion

1st fusion

1st fusion (Welding line)

2nd fusion

1st fusion

Figure 5a). Evolution of the grain size in the HAZ at the intersection in the

case of two GTAW, remelting carried out on the sheet steel IF-TiNb, E=

1kJ/cm with a thickness of t=0.8mm, and 5b). CED-GTAW



Table 5. Ferrite Grain Size Observed in the HAZ or at 950°C (Initial Temperature of the Grain Growth) Depending on the Initial State of Precipitation and the Thermal Gradient for the Steel Grade of IF-Ti,

t=0.7mm Initial State of precipitation Without thermal gradient With thermal gradient

(case of welding) Complete precipitation (as received condition )

10 µm (equiaxed grain)

400 µm (elongated grain)

After preliminary dissolution of the TiC

225 µm (equiaxed grain)

800 µm (elongated grain)

7. TRANSMISSION ELECTRON MICROSCOPY (TEM) STUDIES

A basic TEM study was carried out only on the IF-Ti steel grade samples to observe the precipitates in the microstructures (base metals and weld bead heat affected zone). TEM imaging of thin foils revealed similar precipitate spatial distributions for the IF-Ti steel grades. Bright field transmission electron micrographs show fine aligned precipitates and lightly distributed coarse particles within the grain and inter-phase of the IF-Ti steel. Another study [10] showed that for titanium bearing IF steels, when EDX (energy dispersive X-ray spectrometry) traces are obtained from precipitate particles that indicate the presence of Ti only peaks, these precipitates can be assumed to be TiC or TiN. TiC and TiN and can be distinguished by their shape and size, since TiC particles are fine and more or less spherical, while TiN particles are coarse and cubic in shape.

In fact, our observations indicated that IF steels do not arrive at a state of complete recrystallization. This result is contrary to the literature, where static crystallization is defined by a very low dislocation density. Here, mainly a high density dislocation is observed inside the new sub grains and also in the larger grains surrounding them.

The grain boundary structure was also determined by means of TEM studies, especially to identify the presence of grain boundary ledges (<15 µm for our measurements) in the IFS samples. A representative micrograph in these IFS samples is shown in Figure 6a-e. Grain boundary ledges and a small number of grain boundary dislocations can also be observed in these IFS samples. However, no H precipitates were found in the vicinity of the boundaries of our test samples, as shown in Figure 6b, c. A bright field micrograph taken from base metal, BM (Figure 6a), reveals TiC and TiN precipitates in the as received IF-Ti specimen. Figure 6d shows an example of very fine precipitate distribution in the HAZ, mainly in the surrounding area of the grain boundary.

Additionally, the lower solubility of some particles, such as TiS, TiN, is due to their free energy of formation, ∆Hf, and this means that the more negative ∆Hf is, the less soluble the carbides or nitrides tend to be. In our study, only one possibility appears to be the case. In other words, the particles do not dissolve completely during the weld cycle; however, particle coarsening occurs even with a small amount of precipitates.



Figure 6. Bright field transmission electron micrographs showing aligned precipitates within the base metal, weld bead and heat affected zone:

a) base metal (IF-Ti), b) weld bead/HAZ, v = 2.5 m/s, c) HAZ/weld bead v = 2.5m/s, d) HAZ, v = 2.5m/s, e) weld bead/HAZ, v = 2.5m/s, f) weld bead, v = 2.5m/s

As an alternative approach to Zener, [23,24] have given a more practical equation, which takes into account the distribution of grain size invariably present in practical materials. The difference between the two approaches is given in terms of an equilibrium and non-equilibrium situation. Zener’s approach predicts the grain size after a certain time, whereas Gladman’s approach is based on the time during which the grains are still growing [3,14,15, 23,24]. Our experimental measurements showed that the sizes of the carbonitrides (TiC, TiN…) vary from 6 to 15 nm. Thus, it is very difficult to apply both of these approaches in the case of rapid thermal cycling during the welding of IFS, because they cannot estimate precisely the volume fracture (Vf) of the precipitates. Moreover, they require three dimensional grain sizes. Our measurements of the abnormal grain sizes (500–800 µm), however, are based on the simple two dimensional measurements made wih the pictures taken from an optical microscope [1]. In order to estimate the volume fraction Vf, the empirical equation given by Strid and Easterling [3] can give more realistic results because it gives the two dimensional grain diameter. Thus, abnormal grain growth in IFS during the weld thermal cycle is driven by the thermal gradient and absolutely requires a nucleation stage. It can be deduced that even in the presence of stable precipitates, abnormal grain growth occurs at a certain distance from the fusion zone.



8. CONCLUSIONS

The tests carried out in this study make it possible to separate and quantify the influence of the state of precipitation and the thermal gradient on the grain growth mechanisms in IFS during welding operations:

- The thermal gradient has a dominating influence on the grain growth phenomenon. This aspect is shown in the structure as a “directed” or elongated aspect that is a typical characteristic phenomenon occurring in HAZ. It can be modified by the metallurgy (influence of the composition on the transformation temperatures, etc.) and/or by the welding parameters.

- The initial state of precipitation, particularly the titanium carbides (which are the precipitates most likely to dissolve in a temperature range corresponding precisely to that of the appearance of coarse grains), play an important role in the displacement of the interface (γ α) under the effect of a thermal gradient, which is facilitated even if they are very few in number. Moreover, the effect of the precipitates is reinforced with a thermal gradient.

ACKNOWLEDGMENTS

This project is supported by ARCELOR auto applications. The authors are indebted to the research managers of ARCELOR Research Group for their valuable help in this ongoing project.

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