Microstructural Changes in Brazing Sheet due to Solid-Liquid Interaction

Transcription

1 Microstructural Changes in Brazing Sheet due to Solid-Liquid Interaction Aad Wittebrood

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3 Microstructural Changes in Brazing Sheet due to Solid-Liquid Interaction Aad Wittebrood 2009

4 Cover: Section of a plate oil cooler produced by Behr GmbH, Stuttgart, Germany, Copyright: Corus Technology B.V. ISBN: All rights reserved. No part of the material protected by this copyright may be reproduced in any form or by any means without written permission from the author.

5 Microstructural Changes in Brazing Sheet due to Solid-Liquid Interaction Proefschrift Ter verkrijging van de graad van doctor aan de Technische Universiteit Delft, op gezag van de Rector Magnificus, prof. dr. ir. J.T. Fokkema voorzitter van het College voor Promoties, in het openbaar te verdedigen op woensdag 7 oktober 2009 om uur door Adrianus Jacobus WITTEBROOD geboren te Velsen

6 Dit proefschrift is goedgekeurd door de promotoren: Prof. dr. R. Boom Prof. ir. L. Katgerman Samenstelling promotiecommissie: Rector Magnificus Prof. dr. R. Boom Prof. ir. L. Katgerman Prof. dr. ir. R. Benedictus Prof. dr. ir. J.H.W. de Wit Prof. dr. I.M. Richardson Prof. dr. J.Th.M. De Hosson Dr. K Vieregge Voorzitter Technische Universiteit Delft, promotor Technische Universiteit Delft, promotor Technische Universiteit Delft Technische Universiteit Delft Technische Universiteit Delft Rijksuniversiteit Groningen Aleris Aluminium Koblenz GmbH This research was sponsored by Corus RD&T and Aleris Europe

7 Contents Page Contents V Chapter 1 - Introduction to aluminium brazing sheet 1.1 Introduction Production of brazing sheet Processes Vacuum brazing Controlled Atmosphere Brazing Interaction between the solid core and liquid filler metal Thesis organisation References 9 Chapter 2 - Literature overview 2.1 Introduction A literature overview of the interaction between molten Clad and solid Core alloy Brazing sheet related publications Summary brazing sheet related literature Conclusions References 27 Chapter 3 - Metallurgical changes during brazing 3.1 Introduction Materials and Processing Results Grain structure Precipitate distribution Interaction depth between molten clad and solid core alloy Element distribution Discussion References 49 Chapter 4 - Liquid Film Migration 4.1 Introduction What is Liquid Film Migration? Diffusion from a moving boundary Validation diffusion model Experimental Driving force(s) for Liquid Film Migration Coherency strain Theory Chemical free energy Curvature driven (sub) grain coarsening Grain boundary wetting Reduction of particle/matrix interface energy Surface tension anisotropy 64 V

8 4.6 Discussion Liquid Film Migration Driving forces Conclusions References 68 Chapter 5 - Strain Induced Liquid Film Migration 5.1 Introduction Strain Induced Grain Boundary Migration Strain Induced Liquid Film Migration Grain boundary wetting Experimental setup Grain Boundary Wetting Strain Induced Liquid Film Migration Results Grain boundary wetting by the silicon-rich cladding alloy Grain boundary wetting by Gallium Strain Induced Liquid Film Migration Discussion Grain boundary penetration by silicon from the molten clad alloy Grain boundary penetration by Gallium Strain Induced Liquid Film Migration Conclusions References 96 Chapter 6 - Theoretical considerations of SILFM 6.1 Introduction A simple model to describe SILFM Discussion Alternative models Conclusion References 113 Chapter 7 - The SILFM free material 7.1 Introduction Literature and Patent solutions Through Process Model Results Discussion Conclusion References 124 Summary 127 Curriculum vitae 131 Dankwoord 133 VI

9 Chapter 1 Introduction to aluminium brazing sheet 1.1 Introduction Aluminium brazing sheet is a sandwich material used for the mass production of automotive heat exchangers. The thermal management in a car consists of a number of heat exchangers placed in the front section of a car. Their purpose is to keep the engine cool and the passenger cabin at a comfortable temperature. The heat exchangers are made from aluminium brazing sheet. Figure 1.1 shows the position of the heat exchanger in a modern car. Fig. 1.1: Automotive heat exchangers. Courtesy of Solvay Fluor GmbH 1

10 Basically the sandwich material consists of an aluminium core alloy, typically an AA3XXX alloy (containing Mn) or an AA6XXX alloy (containing Mg and Si) with a clad alloy of the AA4XXX series (containing Si). The core alloy gives the final product the desired properties after brazing. The core alloys are designed in such a way that after the brazing cycle, the condition is reached where the core has its optimum properties. Properties like strength and corrosion resistance are the main engineering parameters. The AA3XXX or AA6XXX alloy core is clad at least on one-side with an AA4XXX alloy but for specific applications where corrosion protection is important, the other side of the core alloy can be clad for example with an AA4XXX, AA1XXX or AA7XXX alloy. The AA4XXX alloy used for brazing sheet has a melting range between 570 C and 610 C while the melting range of a typical AA3XXX core alloy lies above 610 C. This difference in temperature between the two alloys is used to join complex shaped products in one shot. At the brazing temperature, typically around 600 C, the AA4XXX clad alloy is completely molten. Due to capillary forces and surface tension differences, the molten clad alloy will flow to connect adjacent pieces. This complete spreading of the molten clad alloy over the component is related to the origin of the word brazing. AA4XXX AA3XXX Brazing process fillet Fig. 1.2: Schematic and cross-section of as-produced brazing sheet and a brazed joint in an inverted T sample. Generally, in brazing a liquid metal is used to join two metallic or non-metallic pieces. The term Brazing is derived from the old English word braes, meaning to cover with brass [1]. Instead of brass, in modern heat exchangers molten aluminium covers the component. Figure 1.2 shows the essential features of what a brazed joint looks like and figure 1.3 shows a schematic lay-out of a heat exchanger. 2

11 Fig. 1.3: Schematic presentation of the different components of a brazed heat exchanger. In figure 1.3, a cooling fluid flows through a tube-like construction. This fluid takes away the heat from, for example, the engine. This heat has to be dissipated at a different place. Dissipation takes place in the heat exchanger; the heat is transferred form the fluid through the tube to the fins. The fins are cooled by passing air, which takes up the heat from the fin surfaces. Brazing of aluminium can take place under different conditions. However, all these conditions have in common the fact that the ever present tenacious aluminium oxide layer has to be removed and re-oxidation of the surface has to be prevented. McCubbin presented an excellent historical overview of the development of the different brazing processes [2]. Although a number of processes were presented, at present only two of them are of industrial relevance, namely vacuum and controlled atmosphere brazing. Prior to vacuum and controlled atmosphere brazing, which were introduced in the 1960s the dip flux brazing process was used for mass manufacturing of automotive heat exchangers. Flux is introduced to remove the oxide layer allowing the molten clad alloy to come into contact with bare aluminium. During dip flux brazing, the heat exchanger is completely submerged in the flux to facilitate brazing by removal of the surface oxide layer. A major disadvantage of dip flux brazing is the necessity for removal of all the residual flux from the aluminium surfaces after brazing. For local brazing, torch brazing is used in combination with a flux applied at the area to be joined. All dip flux brazing and torch brazing processes are a source of chlorides and it is well known that chlorides and aluminium do not go well together in maintaining corrosion resistance of the aluminium. Any residual chlorides will cause pitting corrosion in aluminium [3]. For this reason, all reactive flux brazed aluminium components have to be thoroughly cleaned before use, leading to a substantial increase of the overall cost. 3

12 1.2 Production of Aluminium Brazing Sheet Most core and clad alloys for brazing are produced by Direct Chill casting. Depending on the required properties, the core alloys can be homogenised or nonhomogenised before the clad alloy is bonded to the scalped 3XXX ingot. The clad alloy is first reduced in thickness by hot rolling and subsequently cut and stretched to assure flatness. The surfaces of clad and core alloy are put together before hot rolling. The first passes of the hot rolling process are responsible for adhesion of the two alloys. Care has to be taken that no rolling oils or dirt of any sort enter between the clad and core alloy since this might cause blisters during subsequent thermal cycling. Hot rolling, cold rolling and annealing steps are used to produce brazing sheet in its final gauge and temper. Figure 1.4 shows a typical processing route for brazing sheet production. melting casting homogenisation scalping hot rolling stretching cutting cladding pre-heating roll bonding hot rolling storage cold rolling batch annealing stretching leveling end annealing Fig. 1.4: Production process brazing sheet. 1.3 Brazing processes Vacuum brazing Vacuum brazing does not require a flux to be applied to the surface of the assembled product to facilitate brazing. To promote vacuum brazing, magnesium is added to the clad alloy [4]. At low-pressure levels (<10-4 mbar) and relatively high temperatures (>500 C), magnesium starts to evaporate from the surface of the clad alloy [5]. The magnesium evaporation follows a burst pattern indicating a disruption of the oxide layer [6]. The evaporated magnesium acts as an atmospheric getter in the vacuum brazing chamber for the remaining oxygen and water vapour. Locally the aluminium surface is protected against re-oxidation. The volatilization of magnesium from the clad alloy disrupts the surface oxide layer, thereby enhancing the clad alloy flow, the molten metal will start to wet the adjacent surfaces [7, 8, 9]. This wetting action is responsible for the joint formation. After a sufficient time at peak temperature, the 4

13 component is allowed to cool down. During cooling the joints will solidify, resulting in a fully brazed or joined component. Molten clad alloy Magnesium vapour Aluminium oxide Al-Si-Mg clad Figure 1.5 shows a schematic presentation of the actual vacuum brazing process. Fig. 1.5: Schematic presentation of the vacuum braze process Controlled Atmosphere Brazing Although vacuum brazing provides products of high quality, the process is discontinuous and demands high investments in equipment moreover maintenance costs are substantial. In the early 1970s this was recognised by ALCAN and a process was developed using a non-corrosive flux in a nitrogen atmosphere. This process was patented [10,11] and came into use in the early 1980s. The process nowadays is called Nocolok brazing being one of the available Controlled Atmosphere Brazing (CAB) processes. The flux has been designed to melt just before the melting of the clad alloy and is capable of dissolving the aluminium oxide layer. When the clad alloy melts it is no longer hindered in its movement by the oxide layer. Figure 1.6 shows schematically how the process works. Molten flux with dissolved oxide and prevents re-oxidation Molten clad alloy Aluminium oxide Flux particles Al-Si clad Fig. 1.6: The Nocolok brazing process. 5

14 The main advantage of this process is the use of a non-corrosive flux. Due to its chemistry (a near eutectic mixture of KAlF 4 -K 2 AlF 6.H 2 O) no chlorides are left on the brazed component and corrosion is not initiated or accelerated by chloride ions. After brazing, the flux remains on the surface as a ceramic layer not influencing any of the properties of the underlying aluminium. The residual flux by itself gives some additional corrosion protection of the brazed aluminium parts during accelerated testing [12,13]. Some conditions have to be met to assure good brazing. The nitrogen needs to have a dew point of below 40 C and oxygen level has to be less than 100 ppm. Magnesium can reduce the efficiency of the flux and to maintain an acceptable level of flux applied in production the Mg content in both clad and core alloy should not exceed 0.3 wt%. Brazing can be done at atmospheric pressure allowing a continuous furnace design with a relatively low capital investment and maintenance costs. At date the vacuum and Nocolok brazing processes are used for 90% of the industrial production of automotive heat exchangers. 1.4 Interaction between the solid core and liquid filler metal All brazing processes have in common that at a certain moment in time in the brazing process; molten aluminium is in contact with solid aluminium. During this stage undesired interactions can take place between a liquid and a solid phase. The interaction between the liquid and solid phase can be divided into three different types. Although the interactions have been given various names, in this thesis the description as given by Woods [14] is used. The types of interactions with their definitions are: 1. Dissolution The static liquid braze metal dissolves the solid from the surfaces with which it is in contact. Figure 1.7 shows that with higher temperatures and longer holding times, the fin and core start to dissolve. The liquid metal from the molten clad alloy is dissolving the fin and tube core alloy, resulting in an increased volume of dendritic brazed joint structure after solidification of the specimen. 6

15 595 C for 5 min. 610 C for 2 min. 100 µm 50 µm 625 C for 2 min. Fig. 1.7: Dissolution of fin and tube alloy. 2. Erosion (Guttering) Liquid metal moves over a surface, rapidly dissolving the solid alloy over which it is flowing. An external force makes the liquid clad alloy flow. During this flow the liquid clad alloy will dissolve the solid alloy over which it is flowing and will carry it away from its position of dissolution, resulting in a reduced thickness. The external force can be gravity or surface tension differences. In figure 1.8 the liquid filler has been running in a narrow crevice by capillary forces, resulting in a circular area within the tube alloy that has been dissolved. Fig. 1.8: Guttering of a tube wall by erosion [14]. 7

16 3. Liquid Film Migration A static reaction during which a solid structure is consumed by a non-convection but moving liquid front, lowering the free energy of the system, leaving behind a solid with a different composition. Figure 1.9 shows the changes caused in the core alloy by what is called Liquid Film Migration (LFM). The liquid clad alloy has penetrated the core alloy leaving in its wake an area free of precipitates. Notice that a dendritic structure as seen in figures 1.7 and 1.8 is not present after the LFM reaction. Precipitate free area Fig. 1.9: Changes caused by LFM [15]. The redistribution and new structure created by the LFM reaction is susceptible to corrosion attack [16] and thereby reduces the lifetime of the brazed component. Also the interaction of the clad alloy with the core alloy reduces the amount of liquid clad alloy capable of forming joints. The objective of this thesis is to understand why and how LFM is occurring and how an alloy system and its processing can be designed in such way as to prevent or minimize the harmful effects caused by LFM. In the industry, the current trend is to down gauge brazing sheet for cost reduction. This will result in that the relative area affected by LFM will increase making even less material available to give the component its strength and corrosion resistance. 1.5 Thesis organisation First the relevant literature will be discussed in relation to liquid film migration observed in metal and non metal systems. A summary will be given on the possible influence of the different parameters on the onset or severity of liquid film migration. 8

17 Chapter 3 deals with the processing of the materials and experiments carried out for this thesis. Chapter 4 relates the measurements and observations to liquid film migration and the diffusion profiles measured are used to calculate the velocity of the liquid film. An inventory is given of the possible driving forces in the studied system. In chapter 5, an explanation is given for the observations and measurements in relation to the recrystallization process known as Strain Induced Boundary Migration. It will be demonstrated that wetting by the liquid clad alloy of the grain boundaries, is a pre-requisite for the onset of liquid film migration. The presence of residual strain is the main driving force for LFM next to the reduction of grain boundaries. The findings of the previous chapters are then compiled and a theoretical model is derived. This model can be used to qualitatively describe the onset and severity of LFM based on process and material parameters. The last chapter makes an attempt to validate a Through Process Model for the use of processing brazing sheet which is less susceptible for LFM. 1.6 References 1. G. Humpston and D.M. Jacobson, Principle of Soldering and Brazing, ASM international, Materials Park, Ohio, USA, 1991, ISBN: J.G. McCubbin, 3 rd International Congress Aluminum Brazing, Düsseldorf, Germany, May 26th- 28th, ASM Handbook, vol 13, Corrosion, ISBN , C.J. Miller, United States Patent 3,321,828, D.K. Creber, J. Ball and D.J. Field, SAE paper W.L. Winterbottom, Welding J., October 1984, W.A. Anderson, Weld. J. Research Supplement, October 1977, s. 8. B. McGurran and M.G. Nicholas, Weld. J. Research Supplement, October 1984, B. McGurran and M.G. Nicholas, J. Mater. Sci. 19 (1984) E.R. Wallace and E.W. Dewing, United Stated Patent 3,951,328, W.E. Cooke, United Stated Patent 3,971,501, P.E. Portin, W.M. Kellermann and F.N. Smith, SAE paper , D.E. Davies and R.M. Prigmore, Proceedings of the 10th ICMC, Madras, India, 1988, p R. Woods, 11th International Invitational Aluminum Brazing Seminar, October 24-26, 2006 Livonia, MI, USA. 15. A.J.Wittebrood, S. Desikan, R. Boom and L. Katgerman, Materials Science Forum Vols (2006) pp

18 16. S. Meijers, Corrosion of aluminium brazing sheet, thesis, Delft University of Technology, ISBN:

19 Chapter 2 Literature overview 2.1 Introduction All brazing processes have in common that at a certain moment in time in the process molten aluminium is in contact with solid aluminium. During this stage interaction takes place between a liquid and a solid phase. From the literature it is known that this interaction can cause changes in the structure and properties of the core alloy involved. These changes are regarded as detrimental to the final product performance. In figure 2.1 two examples of brazing sheets after a brazing cycle are presented, one example reveals hardly any interaction (top right) while the second example (bottom right) shows severe interaction. no interaction Brazing cycle strong interaction Brazing sheet prior brazing line scan area Fig. 2.1: Samples of minimal and severe interaction between clad and core alloy. 11

20 Electron Probe Micro Analysis, (EPMA): of the affected area showed a remarkable redistribution of the alloying elements. As an illustration of this, figure 2.2 shows an area from figure 2.1, which has been analysed by EPMA. The line scan which was made over this area showed a non uniform distribution of the alloying elements between the core and the clad alloy. Fig. 2.2: Line scan area. Table 2.1 shows the chemistries of the core and clad alloy of the pre-brazed sample. Table 2.1: Chemistry of a modified 3003 core and clad alloy AA4045 in wt%. Cu Fe Si Mn Mg Ti Cr Core (mod. 3003) Clad (AA4045) <0.02 < <0.04 <0.04 <0.02 <0.03 Figs. 2.3 and 2.4 show the levels of alloying elements in the affected area of figure 2.2. Fig. 2.3: Element distributions of figure

21 Fig.2.4: Detailed element distributions of figure 2.3. In the affected zone a distinct re-distribution of the alloying elements has occurred. The nature of these changes is not well understood at the present time. This literature overview collates the publications related to the interaction between molten clad and core alloy during brazing. The main focus is on the literature possibly related to Liquid Film Migration as mentioned in chapter 1. This overview will show that there is a limited number of publications available describing the metallurgical processes responsible for the observed changes. 2.2 A literature overview of the interaction between molten Clad and solid Core alloy Brazing sheet related publications Most of the clad alloys contain between 7 and 12.5 wt% of silicon as the main alloying element. The brazing cycle has a peak temperature well above the melting temperature of the clad alloy in use. This means that at a certain stage the solid cladding alloy transforms into a molten metal. This molten metal will interact with its surroundings and especially with the core alloy. In 1943 Miller [1] mentioned that not all of the clad alloy theoretically available contributed to flow due to interaction of the molten clad with the solid core alloy, at a certain temperature the core alloy starts to dissolve in the clad alloy. The exact nature of the interaction was not described and no pictures were presented. However, this article demonstrated clearly that the molten clad and core alloy showed an interaction that could have a detrimental effect on the brazing process. This is illustrated schematically in figure

22 Fig. 2.5: Availability of molten clad alloy [1]. The influence of the effect of silicon diffusion prior to brazing on the clad flow is discussed by Terrill [2]. He demonstrated that the calculated and measured clad flow showed a correlation but with an offset as can be seen in figure 2.6 Fig. 2.6: Theoretically and practically available molten clad alloy [2]. His explanation for the offset was that melting occurs at the grain boundaries and much of the fluid phase is held between grains by capillary forces and thus cannot flow to form fillets. No further details were given. 14

23 Woods and Robinson [3] studied the flow of brazing fillers in dip flux brazing and observed the same offset. They described the interaction responsible for reducing the braze flow due to core solution. Later, the first observations were made [4] relating the amount of cold work as a result of forming prior to brazing and the grain size of the clad alloy on brazeability. Flow in materials that showed insufficient molten clad alloy flow could be partly or completely restored by applying enough cold work to allow full re-crystallisation of the core alloy during the brazing cycle. The critical amount of cold work was considered to be between 3% and 20%. The mechanism behind the improvement of brazeability was not explained; only the restoration of brazeability by applying cold work was mentioned. The chemical composition of the clad alloy has an effect also on the amount of the interaction between the molten clad and core alloy [5]. Higher silicon levels lead to more dissolution of the core material. Addition of copper to the clad alloy promotes penetration even more while the addition of iron up to 1.4% did not show any effect on brazeability. No information was given on the state of the core alloy used in this investigation. One of the first publications reporting an in-depth study of the interaction between molten clad alloy and core alloys came from Schmatz [6]. The interactions between AA6951 and AA3004 core alloys with a 4XXX cladding were studied. In this work, the interaction between the AA6951 core alloy and the high silicon liquid clad alloy was called silicon grain boundary penetration. The liquid originating at the interface between the clad alloy and core alloy becomes enriched in solute which lowers the melting point and promotes penetration along the relatively low melting point grain boundaries. A metallographic analysis of the different areas also showed a significant re-distribution of elements and dispersoids. According to the author, the grain growth appears to be caused by a mechanism similar to that observed during liquid phase sintering. Liquid Phase Sintering (LPS) is a process used to sinter powders to a monolithic structure with the help of a liquid phase between the powder particles [7]. This liquid phase enhances the rate of inter particle bonding during sintering. Grain boundary penetration would have a detrimental effect on the corrosion resistance of the brazing sheet. Table 2.2 shows the chemistries of the alloys studied. Table 2.2: Aluminium Alloy Composition in wt% [6]. Alloy Si Fe Cu Mn Mg Zn Bi

24 Electron microprobe analysis of the penetrated region and the liquid adjacent to the region of alloy AA6951 showed remarkable differences. Two different penetrated areas could be observed. One in which the penetrated area showed no precipitates and copper and magnesium were lower in the penetrated areas compared to the original core levels whilst silicon was found in a concentration at the solubility level in aluminium at the brazing temperature, see figures 2.7a+b. In the second penetrated area, precipitates of a combination of Mg and Si or Mg, Si and Cu were found, see figure 2.8. Fig. 2.7a: Precipitate-free liquid boundary [6]. Fig. 2.7b: Liquid boundary with precipitates[6]. Grain boundary penetration was also observed in the AA3004 alloy; however, to a much lesser extent than with the AA6951 alloy. The liquid boundary with dispersoids from the AA3004 alloy showed a remarkable increase in silicon and manganese, as can be seen in figure 2.8. Again it showed that accumulation of alloying elements took place in the liquid boundary. Fig. 2.8: Liquid boundary with precipitate of alloy 3004 after brazing [6]. Schmatz [6] postulated that the degree of silicon grain boundary penetration is governed by the liquid/solid interfacial energy relationship and the grain boundary energy. The dihedral 16

25 angle (θ ), made by the liquid with the solid at a grain boundary is related to the interfacial and grain boundary energies through equation 2.1. cos θ γgb 2 = 2γ (2.1) sl γ gb = specific grain boundary energy (J/m 2 ) γ sl = solid/liquid interface energy (J/m 2 ) This equation has also been described in several metallurgical textbooks [8,9], whereγ sl is however presented as γsl for a grain boundary/liquid system. Figure 2.9 shows schematically the relationship between the dihedral angle (θ ) and the two energies. γ sl solid (grain) θ grain boundary γ gb γ sl liquid solid (grain) Fig. 2.9: The balance of surface and grain boundary tensions at the intersection of a grain boundary with a free surface. The dihedral angle (θ ) is the angle formed as a result of the different forces being in equilibrium. Gündüz and Hunt [10] used equation 2.1 to determine the grain boundary and solid-liquid surface energy of different systems. In a binary Al-10 wt% Si system γ sl =168.95±21.96 mj/m 2 and γ gb =336.50±47.11 mj/m 2 were measured. Using these numbers in equation 2.1 gives a dihedral angle (θ ) of 2.6 meaning there is excellent wetting by the liquid. Okamoto et al. [11] published work on the brazeability and erosion depth of molten Al-10 wt% Si clad alloy on core alloys made with single additions of iron, nickel, zirconium and silicon. The depth of erosion was measured by brazing non clad cold rolled core alloys with an Al-10 wt% Si clad alloy. The core alloys were given a cold reduction of 80% prior to brazing. The main conclusion was that the depth of erosion depended on the grain size after recrystallization of the core alloy during the brazing cycle, whereby small grains gave more erosion. 17

26 A correlation was found between the recrystallization temperature and erosion depth. Okamoto also suggested that the amount of elements in solid solution plays an important role in the occurrence of the amount erosion. The first recorded attempt to reduce silicon penetration was presented by Engström and Gullman [12]. An AA3003 aluminium inner liner was used between the molten clad and core alloy. The AA3003 inner liner acts as a diffusion barrier and extends the time before silicon from the clad penetrates the core alloy, resulting in less core penetration and improved braze-ability. Yamauchi and Kato [13] also showed that the amount of cold work prior to brazing had an effect on the erosion depth occurring during brazing. The same publication mentioned the effect of dispersoids on the erosion or penetration depth. Several AA3003 and AA3003+Cu alloys were processed in such a way that different dispersoid distributions were created in the core alloy. These dispersoids play an important role in the recrystallization taking place during the brazing cycle. It was stated that a delay of recrystallization before brazing results in more erosion of the core alloy. The number density of small dispersoids (<100nm) that can prevent recrystallization by retarding grain boundaries or sub-grain boundaries from moving and creating a new grain structure [14] play a dominant role in the occurrence of erosion or silicon penetration. The paper did not address the issue of why or how retarding recrystallization of the structure can lead to more erosion. Eichhorn et al. [15] found that the level of iron in the core alloy significantly influenced the silicon penetration into the core alloy. Low levels of iron were found to be beneficial by increasing molten clad alloy flow. However no explanation was given for this observation. Nylén et al. [16] presented a model to explain how silicon core erosion or grain boundary penetration occurred due to the action of grain boundary tensions. The unequal density of grain boundaries on either side of the liquid penetrating film results in a curvature. At a small curvature, this can affect the equilibrium solubility of silicon as described by the Thomson- Freundlich or Gibbs-Thomson equation 2.2. Using this equation together with Fick's first law of diffusion (2.3) an attempt was made to calculate the speed of the advancing liquid film. This number was correlated with microscopic measurements of the rate of advance of the liquid film. Cl 2γVm ln = Thomson-Freundlich or Gibbs-Thomson equation (2.2) C RTd s C l C s = solubility (Si) in the liquid film (mol) = solubility (Si) in the adjacent solid phase (mol) 18

27 γ = specific grain boundary energy (J/m 2 ) V m R T d = molecular volume (m 3 /mol) = universal gas constant (J/molK) = temperature (K) = grain size (m) dc J = D Fick's first law of diffusion (2.3) dx J D dc dx = flux (mol m/s) = diffusion coefficient (m 2 /s) = concentration gradient If 2 γv m << RTd then equation 2.2 can be approximated by equation 2.4. When the following data are used, this assumption appears to be valid. γ = 0.1 N/m V m = 1x10-5 m 3 /mol R T d = J/mol.K = 873 K = 4x10-5 m CV l m2γ C 2 (2.4) RTd Nylén et al. [16] then combined equation 2.3 and 2.4 to give equation 2.5 in order to calculate the speed of the advancing liquid film. J dc C 4CV l mγd = D = D = (2.5) dx t RTdt The following values were used: C = 0.1 l V m = 1x10-5 m 3 /mol γ D = 0.1 N/m = 1x10-7 m 2 /s, diffusion coefficient of Si in the liquid phase 19

28 R T d t = J/mol.K = 873 K = 4x10-5 m = 5x10-6 m (thickness liquid film) A speed of 0.03 µm/s was calculated while a rate of ±0.1 µm/s was deduced from metallography. However, some of the data used are incorrect, the specific grain boundary energy (γ) for aluminium according to [15] is J/m 2 and D has a value of 2.15x10-9 m 2 /s according to Du et al. [17]. Using these numbers, the calculated speed becomes µm/s, which is far less than observed. This would mean that the concentration differences across the liquid film might not be the driving force for the advance of the liquid film into the core. A good example of how small grains and small dispersoids can cause severe brazing problems was presented by Bjørneklett et al. [18]. An extruded alloy strengthened with nanometer sized aluminium nitride particles was brazed with an AlSiMg clad alloy. The grain size of the strengthened alloy was of the order of several microns. Severe penetration of the clad alloy was observed. In the base material the same observation as made by Schmatz [6] was reported, the liquid penetrated the base material and changed the chemical composition of the affected area. At the periphery of the affected area an increase of iron and silicon was observed. In 1997, Woods [19] introduced the term Liquid Film Migration (LFM) for the penetration of the molten clad alloy into the core alloy. A more detailed analysis was done on the affected area and the layer between the affected and non-affected area. Essentially all elements except for Si were removed by a liquid film from the affected area. The level of silicon in the affected area was found to be equal to the maximum soluble amount of silicon in solid aluminium at 600 C being the brazing temperature. All elements were now concentrated in the enriched solidified film. Liquid Film Migration has been observed in many other systems but so far has not been recognised to occur in aluminium brazing sheet. Again it was stated that the amount of cold work prior to brazing had a significant effect on the depth of penetration. Yang and Woods [20] presented at the same conference a number of possible mechanism that could support LFM in aluminium brazing sheet. A relation between cold deformation and the LFM distance after brazing was presented. More attention was paid to possible mechanisms that could support LFM. A number of driving forces were proposed that could supply the energy to explain the kinetics of LFM in brazing sheet. In a subsequent chapter all possible driving forces for the observed changes will be addressed and discussed. 20

29 Wittebrood et al. [21] published the effect of different processing routes applied to a standard brazing sheet alloy. The main conclusions from the investigation were that the homogenization temperature and the sheet gauge of the inter-anneal before the last reduction step had an effect. The homogenization determined the amount of dispersoids and these will have an effect on recrystallization as previously reported by various authors [13,16,20]. The gauge at which inter-annealling is done determines the amount of strain applied before final annealing. This controls the grain size in the final recrystallized condition. Grain size as such also has a significant effect on LFM. It was also found that the amount of manganese in solid solution did not have any effect on the occurrence of LFM. Figure 2.11 shows a correlation between the penetration depth of the liquid film into the core alloy and size of the brazing fillet formed. Due to silicon penetration of the core alloy, the amount of silicon available in the clad alloy is reduced; therefore the amount of clad alloy available to form fillets diminishes which results in smaller fillets. In 2000 Nylén et al. [22] gave a paper on the topic of LFM. Again the mechanism [16] as described by equation (2.6) was used to explain the moving liquid film while no additional data were presented to support this. Once more the effect of recrystallization or its absence on the occurrence of LFM was emphasized. In all previously mentioned publications, LFM was measured either by optical or scanning electron microscopy. Wittebrood et al. [23] used Differential Scanning Calorimetry (DSC) to correlate the penetration depth caused by LFM with the heat input or dissipation during melting and solidifying of the clad alloy. A good correlation was found between the ratio of enthalpy of solidification ( H sol. ) / melting ( H melt. ) and the fillet size as can be seen in figure Fig. 2.11: Correlation between depth of penetration and fillet size [23]. 21

30 During the heat up phase of the brazing cycle, the amount of energy required to melt the clad alloy is represented by H melt.. During the time the molten clad alloy is in contact with the solid core alloy, penetration of the clad alloy into the core will reduce the amount of available molten clad alloy. This will result in a difference between the melting enthalpy of the clad alloy at the beginning of the braze cycle and solidification enthalpy of the clad alloy at the end of the braze cycle. Fig. 2.12: Relationship between ratio of melting and solidification enthalpies to fillet size [23]. DSC proved to be sensitive enough to measure the effect of cold work on the onset of recrystallization during brazing. A reduction of the melting/solidification ratio correlated well with the penetration depth. Figure 2.13 shows the effect of cold work on the ratio solidification/melting. Fig. 2.13: Melting/solidification ratio versus straining with the associated grain structures [23]. 22

31 In 2001 a paper by Yoon et al. [24] on the fabrication of double side clad AA3003 alloy for fin applications appeared. Sagging resistance is the ability of the material to withstand bending under its own weight at elevated (braze) temperatures. The sagging resistance of the material was related to the erosion depth. For fin materials, good sagging resistance is important since fins provide structural integrity to the heat exchanger during brazing. Again the amount of cold work put into the material prior to brazing played an important role in the onset of LFM. Reference was made to the penetration of molten filler into the core along sub-grain boundaries. No explanation was given for the driving force responsible for this penetration. In 2002, the same institute as in reference [24] investigated the effect of the Si level in the core alloys on the depth of penetration [25]. The same findings as described in [13, 20, 22] were reported. Summarizing these, the size and amount of dispersoids are important in influencing the recrystallization behaviour which in turn has an effect on the occurrence of LFM. The amount of Si in the core alloy in the chemistry window investigated determined the formation of fine dispersoids. The production processing sequence has a distinct effect on LFM. The use of continuous annealing compared to batch annealing showed that the worst case of LFM is encountered at different deformation levels for both annealing processes [26]. Continuously annealed material showed much smaller grain sizes that led to more diffusion from the molten filler into the core compared to the coarser grained batch-annealed material. However, continuous annealing is not a way to eliminate LFM, it only shifts the onset of LFM to a different strain level. No explanation was given as to why LFM is occurring. Observation of LFM in a non aluminium brazing systems was made in a publication from McPhee et al. [27]. He described for an Al-Cu-Mg alloy how a moving liquid film could be hindered in its motion by a localised concentration of iron. It was stated that iron has a limited solubility in the solid matrix and the liquid film, with solubility in the liquid film being the highest. Accumulation of iron in the liquid film takes place when it sweeps through the solid matrix. At a certain moment the liquid film is saturated with iron and iron-containing phases will crystallize in the liquid film. When the amount of iron containing phases becomes too high, the particles will not keep up with the moving liquid film due to drag and will therefore become separated from the moving film. Similar features have been observed in brazing sheet. Figure 2.14 shows schematically how liquid pools or islands are being detached in the wake of the moving liquid film. Figure 2.15 is an actual observation in brazing sheet of such detachment. 23

32 Fig. 2.14: Schematic illustration of detachment of a Fe-rich liquid pool from an advancing liquid film.[27] Fig. 2.15: Detachment of a cluster of particles from a moving liquid film Summary brazing sheet related literature In the literature, a number of observations and conclusions regarding the interaction between the clad and core are presented and are often identified by different names including; 1 Core penetration 2 Core erosion 3 Core dissolution 4 Grain boundary penetration 5 Grain boundary melting 6 Liquid Film Migration 7 Braze metal penetration 8 Silicon penetration The microstructure and the chemistry of the core alloy change completely in the zone where interaction takes place. From the literature it seems that the names mentioned are all 24

33 describing the same type of interaction between the molten clad and core alloys during brazing. According to various authors, the interaction depends on the following parameters: Chemistry Alloy chemistry seems to be of importance; especially detrimental are those elements capable of forming dispersoids in aluminium such as Mn, Zr Cr and Ni [13]. Alloys with a relatively low melting point like the 6XXX series are also more susceptible [6, 16]. Whether the low melting temperature or the segregation of elements to the grain boundaries causes this, or the possible smaller grain size of the material, is not fully clear. Homogenization Time and temperature at elevated temperature determine both size and number of the dispersoids formed during homogenization of an AA3XXX alloy [28]. Small particles are known to be responsible for grain boundary pinning (Zener drag) during recrystallization. Grain growing during recrystallization in an alloy containing a dispersion of particles is acted on by two opposing pressures [14]. The net driving pressure for growth P is: P P D P P z C 2 αρgb 3FVγ b 2γb = 2 d R = (2.6) P = total driving pressure P = driving pressure for growth D P = Zener pinning pressure Z P = boundary curvature pressure (Gibbs-Thompson relation) C α = constant of around 0.5 ρ = dislocation density (m -2 ) G = Shear modulus (GPa) b = Burgers vector (m) F = volume fraction particles v γ b = grain boundary energy (J/m 2 ) d = particle diameter r(m) R = grain boundary curvature (m) Grains will only grow if P is positive. When d becomes very small P z will rise. To maintain a positivep, P has to increase. This can only be achieved when the dislocation density ρ D 25

34 is increased. Cold working or deformation increases the dislocation density. This helps to explain why the alloys with small dispersoids [13,18,20,22] are difficult to recrystallize at certain deformation or dislocation levels. Processing Processing and especially annealing plays a role in determining the grain size in the core alloy of the final product. Large strain levels prior to final annealing will result in smaller grain sizes. There is evidence, however, that dislocations introduced by cold rolling can promote precipitation of small dispersoids in an AA3003 alloy [29]. This can then influence the recrystallization behaviour of the core alloy during brazing. Grain size Grain size or grain boundaries play a role in the onset of the interaction [24,27]. No reference was found relating texture to liquid film migration. Grain boundary orientation Only one reference [16] reported whether grain boundary orientation had an influence on the severity of interaction. No relation between mis-orientation and the formation on liquid films could be found. Cold Deformation According to Yamauchi and Kato [14] there is a minimum amount of strain, generally 1 to 5 %, below which recrystallization of the core material will not occur. This is the deformation area in which all published data show severe interaction between the molten clad and core alloy. Little is published on the mechanism describing how a liquid metal interacts with the sub-grain or grain structure at brazing temperature. 2.3 Conclusions There is still a substantial lack of knowledge and understanding regarding the interaction between the brazing sheet cladding and core alloys as described in the literature. Although Liquid Film Migration was mentioned in several publications so far no serious attempt has been made to investigate the wealth of publications on this phenomenon in other alloy systems which may be applicable to aluminium brazing sheet. Publications by for example Yoon [30, 31], Brechet and Purdy [32], Barker and Purdy [33], Kuo and Fournelle [34], Rabkin [35] and many others can provide valuable input to validate the observations made in brazing sheet in the context of LFM. 26

35 In a discussion with Professor Gary R. Purdy from the Materials Science and Engineering department at McMaster University, Hamilton, Canada, being co-author of a number of articles related to liquid film migration [32, 33], he commented that observations made in brazing sheet show the marks of liquid film migration, however this has to be confirmed by additional experiments. If indeed liquid film migration is taking place in aluminium-brazing sheet, the most interesting question is, why is it taking place? The driving force behind liquid film migration is one topic that is still under discussion. Several potential sources are mentioned as the driving force behind liquid film migration such as: 1. Coherency strain 2. Chemical driving force (energy of mixing) 3. Triple point instability (grain boundary wetting) 4. Curvature grain (sub grain) coarsening 5. Reduction of particle/matrix interfacial energy 6. Surface tension anisotropy There is as yet no consensus as to which mechanism prevails. Also combinations thereof cannot be ruled out at the present time. A thorough study of the literature and a design of experiments should give a possible answer to the questions above. 2.4 References 1. M. A. Miller, Weld. J. Research Supplement 22 December 1943, 596-s, 604-s. 2. J.R. Terrill, Weld. J. Research Supplement, May 1966, 202-s, 208-s. 3. R.A. Woods and I.B. Robinson, Weld. J. Research Supplement, October 1974, 440-s, 445-s. 4. P. Sharples, Weld. J., 54 (1975), 164-s, 169-s. 5. I. Okamoto and T. Takemoto, Transaction of JWRI, vol. 10, no 2, 1981, D.J. Schmatz, Weld. J. Research Supplement, October 1983, R.M. German, Liquid Phase Sintering,1985 Plenum press, New York, ISBN D.A. Porter and K.E. Easterling, Phase Transformation in Metals and Alloys, second edition,1991, p.128, ISBN , p R.E. Smallman, Modern Physical Metallurgy, Butterworths, London, Fourth Edition 1985, ISBN , p M. Gündüz and J.D. Hunt, Acta Metall. Vol. 33 N. 9 pp , I. Okamoto, T. Takemoto and K. Uchikawa, Transactions of JWRI Vol.12, no.1, 1983, p H. Engström and L.O. Gullman, Weld. J. Research Supplement, October 1988, S. Yamauchi and K. Kato, Keinkizoku 1991, Vol. 41, no.4, pp

36 14. F.J. Humphreys and M. Hatherly, Pergamon, Recrystallization and Related Annealing Phenomena, 1996, ISBN E.G. Eichhorn, A.C. Scott and J.F. Harris, IMechE 1995, C496/067/95, SAE M. Nylén, U. Gustavsson, B. Hutchinson and A. Örtnäs, Materials Science Forum Vols (1996) pp Y. Du, Y.A. Chang, B. Huang, W. Gong, Z. Jin, H. Xu, Z. Yuan, Y. Liu, Y. He and F. Y. Xie, Mater. Sci. Eng. A363 (2003) B. Bjørneklett, Ø. Grong, L. Isdahl, E. Hellum and V. Sande, Weld. J., March 1996, pp R.A. Woods, SAE paper , H.S. Yang and R.A. Woods, SAE paper , A.J. Wittebrood, R. Benedictus and K. Vieregge, Proceedings of the 6 th International Conference on Aluminum Alloys, Toyohashi, Japan, July 5-10, 1998, pp M. Nylén, U. Gustavsson, B.Hutchinson and Å. Karlsson, Materials Science Forum Vols (2000) pp A.J. Wittebrood, C.J. Kooij and K. Vierregge, Materials Science Forum Vols (2000) pp J.S. Yoon, S.H. Lee. and M.S.Kim, J. Mater. Pro. Technol., 111 (2001) J.S.Rung, M.S. Kim and D. Jung, J. Mater. Pro. Technol., (2002) A. Fukumoto, T. Doko, Proceedings of the 9th International Conference on Aluminum Alloys (2004), pp W.A.G. McPhee, G.B. Schaffer and J. Drennan, Acta Mater. 51 (2003) Y.J. Li and L. Arnberg, Acta Mater. 51 (2003), S.P. Chen, N.C.W. Kuijpers and S. van der Zwaag, Mater. Sci. and Eng., A341 (2003), D.N. Yoon, Annu. Rev. Mater. Sci., : D.Y. Yoon, Materials Reviews, 1995, Vol.40, No.4, Y. Brechet and G.R. Purdy, Scr. Metallur. Vol. 22, 1988, pp S.W. Barker and G.R. Purdy, Acta Mater. Vol. 46, 1998, No. 2, pp M. Kuo and R.A. Fournelle, Acta Metall. Mater. Vol. 39, 1991, No. 11, pp E. Rabkin, Scr. Mater., Vol. 39, 1998, No. 6, pp

37 Chapter 3 Metallurgical changes during brazing 3.1 Introduction The brazing process normally takes place at a temperature where the AA4XXX series alloy is molten. Depending on the silicon content of the 4000 series alloy the brazing temperature is between 585 C and 610 C. In this temperature range the liquid 4000 series is in intimate contact with the solid AA3XXX or AA6XXX series alloy. During this stage where the two phases co-exist, an interaction between the two can take place. This interaction has been a subject of study for over 50 years. Purely based on the amount of AA4XXX series alloy present on the aluminium brazing sheet, it is possible to calculate the sizes of the joints or fillets to be formed during brazing. It was Miller [1] who reported that not all available AA4XXX series metal would end up in a joint. An interaction between the AA4XXX and AA3XXX series alloy consumed some of the liquid 4000 series alloy. Due to the heavy gauges and correspondingly high clad alloy amount available at that time, the interaction was not given much attention. The interaction was described in subsequent papers [2,3] but has not received much scientific attention. The interaction was given different names like core erosion, core dissolution grain boundary dissolution and more. The multitude of names suggests that the interaction is not well understood. Schmatz [4] was the first to make a detailed chemical analysis of the different areas affected by the interaction. His conclusion was that the liquid originating at the clad-core interface progresses into the core alloy as a film, changing the element distribution on its way. The original grains of the core changed due to the liquid film passage to large grains. Okamoto and Kato [5] and Yamauchi et al. [6] described the effect of dispersoid forming elements on 29

38 the occurrence of the interaction and concluded that recrystallization inhibitors had a highly detrimental effect. Nylén et al. [7] tried to use the Thomson-Freundlich or Gibs-Thomson equation together with Fick's first law of diffusion to calculate the speed of the advancing film. However, the calculated velocity was only 2% of the actually velocity observed. In 1997 Woods [8] and Yang and Woods [9] published a detailed analysis of the observed interaction. Woods was the first to coin the term Liquid Film Migration. The experimental observation made was the complete redistribution of elements when a liquid film has passed through. In the study by Yang and Woods [9] the amount of cold work applied before brazing was again shown to influence the severity of the interaction. Although some authors presented descriptions of the metallurgical changes caused by the interaction between the solid core and molten clad alloy, the available data and explanations are limited. The scope of this chapter is to present a detailed analysis of the changes caused by the interaction. The following changes have been studied: 1. Grain structure 2. Precipitate distribution 3. Interaction depth between the cladding and core alloys 4. Element distribution 3.2 Materials and processing For the present study three different compositions were cast; two core alloys and one cladding alloy. The chemical composition of the alloys can be found in table 3.1. The composition of the cladding alloy falls within the chemistry window of AA4045. The copperfree 3XXX series alloy is essentially an AA3003 alloy while the copper-containing core alloy does not fall within any Aluminium Association designation for 3XXX series alloys. A 25 kg ingot of the different chemistries was cast in a torpedo shaped mould with water-cooling to give a solidification structure similar to that encountered in the production of industrial DC casting. Blocks measuring 100x80x80 mm were cut for rolling from the as-cast ingpts. Table 3.1: Chemical composition in wt% of the two core alloys and clad alloy used in this study. Si Mn Cu Fe Al Core < Balance Core Balance Cladding 10.2 <0.01 < Balance Before casting 1g of Al/TiB-5/1 of grain refiner for each kg of aluminium was added. 30

39 The core alloy blocks were given different homogenisation treatments (see table 3.2) while the clad alloy was hot rolled at 490 C to 10 mm thickness without homogenisation. The 10 mm hot rolled clad material was machined to a thickness of 5 mm. This thickness was used to clad the core alloy blocks by roll bonding. The homogenised core alloy blocks were prepared as illustrated in figure 3.1. Every core alloy sample was processed in both the clad and non-clad conditions. Fig. 3.1: Dimensions of the processed core and clad/core combinations. Hot and cold rolling of the clad and non-clad blocks were carried out as follows: Hot rolling: Heat-up at a rate of 30 C/hour to 430 C, hold for 3 hrs and then hot rolled down to 4 mm. After hot rolling the rolled sheets were allow to cool naturally to room temperature. To make sure that the material was fully recrystallized after hot rolling, the samples were given an additional heat treatment at 400 C for 1 hour. Cold rolling: The annealed 4 mm material was rolled down to 0.4 mm followed by a final anneal with a heating rate of 35 C/hour to 400 C, a hold of 2 hours and cooling rate of 35 C/hour down to 50 C. Table 3.2: Applied homogenisation treatments. Heat up rate C/hour T hold C t Hold hour Cooling rate C/hour T end C Cycle Cycle Cycle Cycle Cycle Prior to brazing, the brazing sheet samples were stretched 0, 1, 3, 5, 7 and 10% in the plane strain mode to introduce dislocations. The sample width ensured that stretching occurred under plane strain condition. These levels of stretching correspond to strains of 0, 0.009, 31

40 0.029, 0.049, and This stretching simulated the different forming strains that are experienced during the manufacture of heat exchanger components. The actual elongations of each sample were measured. All measurements indicated that the actual stretching was close to the desired level with one exception. The non-homogenised Cu containing core alloy could not be stretched beyond 6.6%. At this level the sample started to crack, meaning the material was at its tensile strength. For convenience in the descriptions and discussions throughout this thesis, the levels of stretching of 0 up to 10% will be used except for the calculations made where a more accurate input is required. Brazing was done in a batch Controlled Atmosphere Brazing furnace under a nitrogen atmosphere using a simulated brazing cycle with a maximum temperature of 595 C and a holding time of 3 minutes. To study the effect of heating rate and holding time, samples were immersed in a salt bath set at 595 C for specific times, after which they were quenched in water. Conventional Light Microscopy (LIM) and Scanning Electron Microscopy (SEM) were used to study the grain structure of non-brazed and brazed specimens. Standard metallographic techniques were used to prepare the samples for LIM and SEM. Two methods were used to reveal the precipitates after brazing: 1. An immersion in 0.5% HF for between 30 and 60 seconds. 2. An immersion in a mixture of 15 g of molybdenic acid, 10 ml of HF and 90 ml of water for 30 seconds. 3. The grain structures were made visible by applying Barkers etch with 4% HBF 4 and 20V DC for 90 seconds. The local chemical composition was measured by using a micro probe (CAMECA SX 100) on polished samples. SEM was carried out with a LEO 438 VP. 3.3 Results Grain structure Figure 3.2 contains the micrographs of all clad variants of the different homogenised core alloys. The micrographs show the precipitates of the core alloy which was clad and the grain structures of both the clad and non-clad core alloys in the as-rolled condition. By comparing the grain structures of each clad and unclad core alloy it is possible to make an assessment of whether a cladding alloy has an effect on the grain structure of the core alloy during processing. The effect of copper can also be assessed since Core 2 contains 0.4 wt% Cu while Core 1 contains none. 32

41 Cycle 1: Non Homogenised- No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Cycle 2: Homogenised 24 hrs at 500 C No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Cycle 3: Homogenised 24 hrs at 550 C No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Cycle 4: Homogenised 24 hrs at 600 C No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Cycle 5: Homogenised 168 hrs at 600 C No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Micrographs taken with the same magnification, sample area(550x350µm) Fig. 3.2: Grain structures of the clad and non-clad differently homogenised core alloys in the as-fabricated condition. 33

42 The core alloys with copper seem to have a smaller average grain size than the non-cu containing alloys. Post-braze, grain structure micrographs were also taken of all the different homogenised core alloys in the clad and unclad conditions and with the different levels of stretching. Cycle 1: Non Homogenised S No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Micrographs taken with the same magnification, sample area(550x350µm) Fig.3.3: Grain structures of core alloys 1 and 2, non-homogenised, after brazing and stretched prior to brazing. 34

43 There is not much difference between all grain structures in this series, only the clad non Cu containing alloy at 10% stretch level shows an increase in grain size. The clad brazed alloys show a different grain structure. In some cases large grains are observed where the clad alloy was located. Especially the 10% stretched Cu containing alloy shows a large grain on top at the place of the original clad alloy. Cycle 2: Homogenisation 24 hrs at 500 C S No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Micrographs taken with the same magnification, sample area(550x350µm) Fig.3.4: Grain structures of core alloys 1 and 2, homogenised at 500 C for 24 hours, after brazing, and stretched prior to brazing. 35

44 Again a difference in grain structure can be seen between the Cu and non-cu core alloys. The Cu core alloys show a significant increase in grain size at 7% stretching while this only becomes noticeable at a 10% stretch level for the non-cu core alloys. Cycle3: Homogenisation 24 hrs at 550 C S No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Micrographs taken with the same magnification, sample area(550x350µm) Fig.3.5: Grain structures of core alloys 1and 2, homogenised at 550 C for 24 hours, after brazing, and stretched prior to brazing. 36

45 The initial grain structure for the Cu core alloys appears to be more uniform. With this homogenisation practice the Cu core alloys change grain size at 5% stretching, while the non-cu core alloys show some change at 7% stretching. Cycle 4: Homogenisation 24 hrs at 600 C S No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Micrographs taken with the same magnification, sample area(550x350µm) Fig.3.6: Grain structures of core alloys 1 and 2, homogenised at 600 C for 24 hours, after brazing, and stretched prior to brazing. 37

46 Here the Cu core alloys change in grain size at 5% stretching while the non-cu core alloys change at 7%. Cycle 5: Homogenisation 168hrs at 600 C S No Cu-unclad 0.4 wt% Cu-unclad No Cu-clad 0.4 wt% Cu-clad Micrographs taken with the same magnification, sample area(550x350µm) Fig.3.7: Grain structures of core alloys 1and 2, homogenised at 600 C for 168 hours, after brazing, and stretched prior to brazing. At 3% stretching the grain size of the Cu alloys increases to a size similar to the thickness of the specimen. The non-cu core alloy shows the same change only at 5% stretching. 38

47 Overall the trend is that the grain structure of the Cu containing core alloy changes at lower levels of stretching when homogenisation temperatures become higher and longer. Table 3.2 shows for all samples used in this study when a change in grain structure has been observed. Whether the grain structure changes are caused by recrystallization or a grain structure altering process will be discussed in a subsequent chapter. With longer homogenisation times, the alloys show a tendency to change micro-structure at lower levels of stretching. Table 3.2: Change in grain structure: N= no recrystallization, Y= recrystallization Strain (stretch %) 0 (0%) (1%) (3%) (5%) (7%) (10%) Core Core Core Core Core Core NH N N N N N N N N N N Y N 24 hrs 500 C N N N N N N N N N Y Y Y 24 hrs 550 C N N N N N N N Y N Y Y Y 24 hrs 600 C N N N N N N N Y Y Y Y Y 168 hrs 600 C N N N N N Y Y Y Y Y Y Y 39

48 3.3.2 Precipitate distribution All samples were re-polished and etched in HF to reveal precipitates and constituents. The next figure shows the two extremes of the interaction between the clad and core alloy of all series from figures 3.3 and 3.7. Cu in core alloy Cycle 1: Non Homogenised Cycle 5: Homogenisation 168hrs at 600 C Fig 3.9: Cross sections of post brazed clad Cu containing core alloys homogenised according to Cycles 1 and 5. 40

49 The non-homogenised, Cu containing, clad core alloy shows a change in precipitate distribution at the surface at 3% stretching. This change becomes larger at higher levels of stretching. The 10% stretched sample shows a change up to 150 µm deep into the original thickness of the core material. The changed area seems to have become free from precipitates. The same alloy homogenised according to cycle 5, shows relatively little change in precipitate distribution. Only the 3% stretched sample shows a moderate change at the surface. Figure 3.10 below shows an enlargement of the 10% stretched Cu cycle 1 core alloy of figure 3.9. Fig. 3.10: Details of the interface between the original core and the affected or changed core. Part of the original core has changed; the affected area in figure 3.10 appears to be free from precipitates. All precipitates seem to be concentrated at the interface and/or in clusters just attached to the precipitate filled interface Interaction depth between molten clad and solid core alloy The remaining original core thickness of the samples used in this study was measured and expressed as a percentage of the original core thickness before brazing. The actual measurements of the original core were compensated for the degree of thinning caused by 41

50 the amount of stretching. The figures 3.11 and 3.12 show the thickness of the two remaining cores for all homogenisations and stretch levels applied. 110 Remaining core thickness (%) NH 24h 500 C 24h 550 C 24h 600 C 168h 600 C Elongation prior brazing (%) Fig. 3.11: Remaining core thickness of core alloy 1 (no Cu). 110 Remaining core thickness (%) NH 24h 500 C 24h 550 C 24h 600 C 168h 600 C Elongation prior brazing (%) Fig. 3.12: Remaining core thickness of core alloy 2 (Cu). As can be seen in figure 3.10, the interaction interface does not appear as a straight line; therefore the scatter in figures 3.11 and 3.12 can be as large as 10%. However, in most cases the affected area grows to a maximum after which the affected area reduces in size, 42

51 meaning the affected area is reduced at higher levels of stretching. The point at which the affected area is reduced correlates with the point at which the core shows a change in grain structure as given in table 3.2. This is especially evident for core alloy 2 (figure.3.12). The non-homogenised core alloy 2 stretched 10% (reality 6.6%) was salt bath treated at 595 C for various times up to 720 seconds. Again, the depth of the maximum interface penetration into the core was expressed as a percentage of the original core alloy thickness prior to the salt bath treatment. The results are presented in the figure Fig. 3.13: Remaining core thickness of core alloy 2 (Cu) Element distribution The figures 3.14 and 3.15 show the SEM image of brazed, non-homogenised core alloy 2 stretched 10% before brazing. An element distribution trace has been made along the line indicated as the left hand picture Interface Interface Original core Affected area Wt% Affected area Original core Si Mn Fe Cu Al Distance (µm) Fig. 3.14: SEM image and line scan of core alloy 2. 43

52 2.5 Original core Interface Affected area Wt% Affected area Interface Original core Si Mn Fe Cu Al Distance (µm) Fig. 3.15: SEM image and line scan of core alloy 2. The figures show that the affected area is essential free from precipitates or constituents. All major alloying elements are concentrated at the interface as are the precipitates and constituents. Since manganese is a slow diffusing element compared to silicon and copper, it can be used as an indicator of were the affected area stops and the original core alloy starts. The measured concentration of the different elements in the affected area and the concentration of the elements as calculated thermodynamically by the software program JMatPro-4.0 are compared in figure Wt% Affected area Si av= 1.18 wt% Mn av=0.17 wt% Cu av=0.08 wt% Fe av < 0.01 wt% Distance (µm) Si Mn Fe Cu Al Si =1.15wt% Mn=0.11 wt% Cu=0.02 wt% Fe< 0.01 wt% Fig. 3.16: Measured at 595 C and calculated concentration of elements in solid solution. Table 3.3 summarizes the measured and calculated concentration in wt% at 595 C. Table 3.3: Measured and calculated concentration in wt% in the affected area at 595 C Si Mn Cu Fe Measured <0.01 Calculated <

53 The measured and calculated values at 595 C can be considered in agreement taking into account the inaccuracies of the line scan measurement. This affected area in which the concentration of the elements in solid solution equals the equilibrium solidus concentrations at the temperature of the experiment is considered to be one of the distinct metallurgical features related to Liquid Film Migration [10]. Another difference in the diffusion profiles can be observed in the following measurements. The first measurement was done at the interface of the non-homogenised core alloy 2 stretched 10% and brazed in a salt bath for 480 seconds at 595 C. The second measurement is from core alloy 2, homogenised for 168 hours at 600 C then stretched 10% and also brazed in a salt bath for 480 seconds at 595 C. The only basic difference between the two samples is that one is non-homogenised while the other is homogenised [Si] (wt%) Non homo Homo Calc diffusion depth (µm) Fig. 3.17: Measured and calculated concentrations of Si in wt% at 595 C. The two measured profiles are remarkably different. For comparison, the bulk diffusion profile for Si has been calculated and put in the graph of fig The diffusion profile of the homogenised material is almost identical to the calculated profile, indicating that only bulk diffusion took place. After brazing, the homogenised core was recrystallized while the nonhomogenised core did not show recrystallization. The difference in profile could not be explained by another form of diffusion like grain boundary or dislocation diffusion. Both diffusion processes are faster compared to bulk diffusion and would result in deeper penetration of the silicon. In a next chapter the difference between the two profiles will be explained. The formation and composition of the different precipitates and constituents have been studied in detail by Marshall et al. [11] and later by Avramovic-Cigara et al [12]. 45

54 A key factor in both studies was the use of a relative both low iron (<0.5 wt%) and low silicon (<0.4 wt%) core alloy. In the publication by Braker and Purdy [11] another key feature is that a non-homogenised core alloy was used. Their analysis showed that before brazing, manganese in the core alloy is distributed between dispersoids and solid solution. The dispersoids are characterised as mainly MnAl 6 and some α -AlMnSi phase and are below 100 nm in size. Predominately, the dispersoids are MnAl 6 due to the low iron and silicon levels in the core alloy. The coarser particles are considered to be orthorhombic (FeMn)Al 6. The amount of manganese in solid solution before brazing can be up to 0.75 wt% which means the solid solution is supersaturated since for the investigated core 2 alloy the theoretical solubility of manganese is below 0.01 wt%. When the alloys are exposed to a brazing cycle which goes up to 595 C, silicon starts to diffuse from the clad alloy into the core alloy. Silicon will destabilize the super saturated manganese solution and force the manganese to precipitate as α -AlMnSi precipitates [13]. In the diffusion zone a Dense Band of Precipitates is formed [11]. By using the correct etching, this band can be made visible as can be seen in figure Fig. 3.18: Band of Dense Precipitates made visible by etching with HF Molybdenum acid. The depletion of manganese in solid solution as a result of the formation of the Dense Band of Precipitates is considered to be responsible for the enhanced corrosion resistance of aluminium brazing sheet [11,14,15]. In the particular case of figure 3.18, the precipitates in the DBP are shown together with an area where the precipitates have been consumed by a possibly moving boundary which has penetrated into the core alloy. Figure 3.19 shows in detail the different precipitates and other particles at the interface between the affected area and the original core alloy. The analysis of the different areas, constituents and precipitates are in line with observation made by Marshall et al. [11,16]. It seems that the liquid film sweeps the precipitates and constituents present in the core alloy and concentrates them at the interface between the affected area and original core alloy. 46

55 According to thermodynamic calculations the constituents and precipitates do not dissolve in the liquid film. JMatPro-4.0 calculates a maximum liquid solubility in the interface for Si of wt%, Mn of 0.57 wt% and for Fe of 0.12 wt% in the Cu free alloy. When Cu is added to the core alloy, it will not change the solubility of the other elements. Fig. 3.19: SEM-EDX analysis of precipitates and constituents in and around the interface. The interface between the affected area and original core alloy shows a high concentration of constituents and precipitates. The liquid is located at the interface and originates from the clad alloy. During the movement of the liquid film it is questionable if the precipitates and constituents dissolve and precipitate during this movement. It is assumed that the interface is composed of a mixture of the clad alloy and core alloy, thermodynamics can be applied to calculate if the precipitates and constituents already present will dissolve and re-precipitate. At 595 C the concentration of the elements in the liquid would be 10 wt% for Si, 0.4 wt% for Cu, 0.56 wt% for Mn and 0.12 wt% for Fe. Norgren et al. [17] showed that in a 0.8 wt% Mn based core alloy, the amount of Mn in solid solution in front of the silicon interface is 0.5 wt%. The excess Mn has precipitated as described by the DBP. So most likely the precipitates already present in the core alloy will not dissolve when a liquid film passes through. After cooling of the samples the dissolved Mn will precipitate since the maximum solid solubility of Mn is at 595 C and is 0.11 wt%. 47

56 At the interface three different groups of precipitates are located. The first group are formed during processing, in this case mainly MnAl 6, then the second group is formed due to the silicon diffusion during the brazing cycle and are mainly α -AlMnSi precipitates and the third group is formed after brazing during solidification of the liquid interface of the sample when cooled down. From this analysis it cannot be determined whether the chemistry of any precipitates or constituents changes composition when in contact with a liquid film containing a high concentration of dissolved silicon. The size of the precipitates and constituents does not seem to change. The observation that the interface is enriched with precipitates and constituents from the core alloy is considered to be additional evidence that the changes observed are caused by liquid film migration [18]. To illustrate even more clearly how the elements are distributed an elemental map has been made at the interface. Fig. 3.20: SEM image and element mapping at the interface. All major alloying elements are present at increased concentrations at the interface. Only silicon is present in solid solution in the affected area. In front of the interface into the core alloy, there is a silicon and copper concentration profile. In figure 3.15 the silicon profile runs from 50 to 120 µm going from high (1.2 wt%) to low content (0.1 wt%) while copper shows a profile from 50 to 100µm going from low (0.1 wt%) to high content (0.4 wt%). 48

57 3.4 Discussion Only few detailed element analyses have been made of the interface and its surroundings. Acoff and Thompson [19] studied the element distributions at the interface of a nickel base alloy 718. They observed an enrichment of the alloying elements at the interface and found that the concentration of the elements in the affected area was at equilibrium concentration at the temperature of the experiment. The results from the present study are in line with this analysis, as were those reported by Woods [8]. Although there is significant circumstantial evidence that the observed changes in Al-Si brazing systems are due to the existence of liquid film migration, conclusive evidence is still lacking. Strong evidence for liquid film migration would make it possible to study the driving force for the movement of the liquid. This driving force could then be correlated with any relevant microstructural and processing parameters that can help to control or minimise the occurrence of liquid film migration. 3.5 References 1. M. A. Miller, Weld. J. 22 (12) 1943, Res. Suppl., pp s. 2. R.A. Woods and I.B. Robinson, Weld. J. 53 (10) 1974, Res. Suppl., pp s. 3. I. Okamoto and T. Takemoto, Transaction of JWRI, Vol. 10, 1981,no 2, pp D.J. Schmatz, Weld. J. (10) 1983, Res. Suppl., pp.267s, 271s. 5. I. Okamoto, I.Takemoto and K. Uchikawa, Transactions of JWRI, Vol.12, (1983), no.1, pp S. Yamauchi, K. Kato and Keinkizoku, Vol. 41, (1991), no.4, pp M. Nylén, U. Gustavsson, B. Hutchinson and A. Örtnäs, Materials Science Forum Vols , (1996), pp R.A. Woods, SAE paper , H. S. Yang and R. A. Woods, SAE paper , S.W. Barker and G.R. Purdy, Acta Mater. Vol. 46, 1998, no.2 pp G.J. Marshall, R.K. Bolingbroke and A. Gray, Metall. Trans B, Vol. 24A, September 1993, G. Avromovic-Cigara, S. Thorpe and B. Cheadle, Prakt. Metallorgr. 34 (1997) 2, Y.J. Li and L. Arnberg, Acta Mater. 51 (2003) S. Meijers, Corrosion of aluminum brazing sheet, thesis, Delft University of Technology, ISBN: S. Tierce, N. Pébère, C. Blanc, C. Casenave, G. Mankowski and H. Robidou, Electrochimica Acta 52 (2006) G.J. Marshall, A.J.E. Flemming and A.Gray, SAE paper , C496/030/95. 49

58 17. S. Norgren, A. Oskarsson and R. Woods, 9th International Invitational Aluminum Brazing Seminar, October 26-28, W.A.G. McPhee, G.B. Schaffer and J. Drennan, Acta Mater. 51 (2003) V.L. Acoff and R.G. Thompson, IBM J. Res. Develop. Vol. 44, September 2000, No. 5,

59 Chapter 4 Liquid Film Migration 4.1 Introduction Yoon and Huppmann [1] were the first to report about a change in element distribution in pure solid tungsten grains when put in contact with pure liquid nickel. During liquid phase sintering the observed changes seem to be caused by a liquid film, first dissolving the tungsten followed by re-precipitation of a nickel saturated tungsten solid solution. The observed effect is presented in figure 4.1. Pepe and Savage [2] reported earlier about the movement of grain boundaries attributed to liquation of titanium sulphide inclusions in 18Ni Maraging steel. No detailed analysis was given, but based on the change of colour after etching for microscopy the authors claimed a change in chemistry caused by the movement of a liquid film. Fig. 4.1: The distribution of nickel in a pair of sintered tungsten spheres [3]. 51

60 The re-precipitated solid solution is the equilibrium solidus composition belonging to the temperature of the experiment. The phenomenon involving the movement of a liquid film and the compositional change caused by the liquid film has been named Liquid Film Migration (LFM). A change in chemical composition in a Fe-Zn alloy system caused by grain boundary motion induced by grain boundary diffusion was observed by Hillert and Purdy [4]. The movement and change in composition was called Chemically Induced Grain Boundary Migration (CIGBM). Both LFM and CIGBM have in common a moving interface, therefore both processes are in generic terms called Chemically Induced Interface Migration (CIIM) [5]. Since aluminium brazing takes place at a temperature at which the clad alloy is at least partly molten, the main focus of this chapter will be to validate the assumption that the observed changes as described in chapter 3 are caused by Liquid Film Migration. 4.2 What is Liquid Film Migration? Liquid Film Migration has been observed in a number of metallic and non-metallic systems. In 1995 Yoon [6] made an overview of the different systems where either Chemically Induced Grain Boundary Migration or Liquid Film Migration was observed. In the meantime a number of other systems like Cu-Ti, Al-Cu [7], Al-Si [8], Ni-Base Alloy 718 [9], Al-Zn-Mg-Cu [10], AA2618 [11], Al-Cu-Mg [12], Ni-Fe-W [13], bischofite [14] and Cu-Sb [15] could be added to the list. In all systems a separate liquid phase is locally formed in a solid phase. This liquid phase starts to move into the solid phase leaving in the wake of the moving liquid phase a composition that is the equilibrium solidus composition at the temperature of the experiment [16]. The occurrence of LFM can be beneficial in the case of liquid phase sintering in powder metallurgy, resulting in less porosity [17]. In welding the occurrence of LFM can lead to cracking [18,19]. In aluminium brazing, LFM can decrease the corrosion resistance [8], compromise formability [20] and brazeabilty [21]. Brazeability is used to describe the ability of the liquid filler to form joints, whereby joint formation is a balance between diffusion, wetting, capillary forces and viscosity. Part of this chapter will be dedicated to provide evidence that the existence of a liquid film is responsible for the observed changes. About the exact nature of the driving force for LFM, there are still some questions. In section 4.5 the main driving forces will be highlighted and discussed with respect to their applicability to the observations made in aluminium brazing sheet. 52

61 4.3 Diffusion from a moving boundary Diffusion originating from a moving interface can be represented schematically the figure 4.2. Affected area Interface v c Core alloy D lattice v Fig. 4.2: Concentration profile around a moving interface during LFM. Diffusion with respect to a frame at the interface migrating at velocity v must satisfy: J = d d x D lattice dc vc =0 dx (4.1) This has the following solution [22]: v D x lattice c= c0 e (4.2) Equation 4.2 is used to verify the existence of a liquid film responsible for the observed changes as described in chapter Validation of the diffusion model Experimental Samples of clad non-homogenised core alloy 2, stretched 10% prior to brazing were immersed in a salt bath at 595 C for various intervals. The time for the samples to reach the target temperature was estimated to be 4 seconds. In addition to this time the samples were held between 10 and 720 seconds at 595 C before they were extracted from the salt bath and quenched in water. After quenching the samples were prepared for line scan measurement at the interface between the affected area and original core alloy as presented in figure 3.13a+b. 53

62 The line scan measurements for silicon were used as input for the validation of equation 4.2. The diffusion profiles were measured singly, triplicate or quadruplicate. In figures 4.3 and 4.4 it is shown how the line scans were used to extract the diffusion data. Diffusion profiles 120 s at 595 C Si (wt%) Av Distance (µm) Fig. 4.3: 4 diffusion profiles and averaged. Detail diffusion profiles 120 s at 595 C Si (wt%) y = e x R 2 = y = e x R 2 = y = e x R 2 = y = e x R 2 = Av Distance (µm) Fig. 4.4: Diffusion profile and correlation. The diffusion profile in front of the interface extending into the core alloy was extracted and correlated with an exponential curve with the basic formula. y Bx = A e (4.3) where v B= (4.4) D lattice 54

63 For all measurements, the diffusion profile was fitted as described above. The exponent of each fit for every time at the target temperature was used to calculate the speed of progression of the interface at that time. The diffusion coefficient used for silicon in aluminium at 595 C was calculated from the data by Du et al [23], giving the valued at 595 C = x m 2 /s. Si Table 4.1: Calculated speed of the interface. Time at 595 C (s) v interface (µm/s) S dev. (µm/s) (s) (s) (s) (q) (q) (q) (t) (t) 0.01 (s) single, (t) triple, (q) quadruple measurement The variation of the velocity of interface with time is represented by figure Interface velocity (µm/s) time at 595 C (s) Fig. 4.5: Interface velocity/ time relationship based on diffusion depth measurements. The data points for 100 seconds and longer have been used to calculate the fit between time and speed. Integration of the formula providing the best fit provides the relation between time and travelled distance. The best fit of the curve between 100 and 720 seconds from figure 4.5 is represented by: 55

64 v = 1.26 t, r 2 =0.97 Integration of this equation results in: s= 2.55 t Figure 4.6 presents the calculated curve and the measured penetration depth of the liquid film Distance (µm) Measured Calculated Time at 595 C (s) Fig. 4.6: Measured and calculated penetration depth. Both graphs are in the same range but the measured data show some deviations at the shorter times. Close examination of the samples, however, shows that at the shorter immersion times a clear interface as described in chapter 3, has not yet developed. This can be seen in figure 4.7 a+b, which shows samples held in the salt bath for 40 seconds at 595 C. In figure 4.7a, no distinct interface has developed, while in figure 4.7b a partial interface has developed. Consequently this makes definition of the exact interface position difficult. Fig. 4.7a+b: Built up interface after 40s at 595 C. 56

65 If it is assumed that the development of an interface as described in chapter 3 is a microstructural feature related to LFM, then the first changes occurring in less than 100 seconds at 595 C are necessary to create conditions for the onset of LFM. In short the liquid clad and core have not interacted sufficiently with each other to show LFM. Yang [8] also mentioned an incubation time before LFM was visible in his experiments. 4.5 Driving force(s) for Liquid Film Migration In the literature, a number of sources of the driving force for LFM are reported. Some of the theories are supported by experimental data, while others are not. Published theories or explanations are: 1. Coherency strain theory 2. Chemical driving force (energy of mixing) 3. Grain boundary wetting 4. Curvature grain (sub grain) coarsening 5. Reduction of particle/matrix interfacial energy 6. Surface tension anisotropy These theories will be discussed and will be assessed for their relevance to the onset of LFM Coherency strain Theory The coherency strain theory can best be explained by figure 4.8. Fig. 4.8 a,b,c: The schematic diagram of liquid film migration. 57

66 According to the coherency strain theory which has been proposed by Sulonen [24] for discontinuous precipitation, the following sequence is causing the formation and migration of a liquid film. What follows is a generic description of the processes taking place. First the system is brought to the temperature where the Al-Si alloy starts to melt. The liquid will penetrate between two aluminium grains in the core alloy as shown in fig. 4.8a. Due to the concentration difference between the liquid and the grains of the solid, diffusion of Si into the solid matrix will take place, see fig. 4.8b. The Si diffusion ahead of the liquid film will cause strain in the grain. The relaxation of the built up strain is the driving force for LFM. The orientation of the grains adjacent to the liquid film will determine the migration direction; see fig. 4.8c. The liquid film starts to move into the grain with the highest level of strain leaving behind a relaxed solid phase. In short, the solute misfit is causing strain in the matrix which provides the energy for making a boundary move. The coherency strain theory is applicable to explain in general terms the movement of any interface which is seen as chemically induced interface migration [5]. The coherency strain theory has been treated mathematically first by Hillert and Purdy [4,25] and later by Yoon [5,6] who published a critical validation of coherency strain theory in a liquid phase sintered 85Mo-15Ni system. The theory allows the calculation of the strain energy developed per unit volume in a matrix into which solute atoms have diffused. The strain energy can be described by the following formula for the aluminium brazing system. G= 2 E η 1 υ l 0 ( X ) 2 Si X Si (4.5) Where E is the modulus of elasticity, υ is Poisson s ratio; η is the atomic misfit parameter, l X Si is the mole fraction of silicon at the strained diffusion interface which according to the right hand graph of figure 3.15 is 0.012, while the diffusion Interface and is X Si is the mole fraction of silicon far away from The observed temperature of the experiment was 595 C, hence the modulus of elasticity and the Poisson ratio need to be calculated for this temperature for aluminium. Since the strain developed is dependent on the orientation of the grain, the appropriate E modulus for a given orientation has to be used. To calculate the elastic constant values for the 100 and 111 directions for aluminium, data from Sutton [26] was used and extrapolated to 595 C. E 100 and E 111 have values of respectively 39.3 and 51.7 GPa and υ 100 and υ 111 are respectively around 0.31 and The atomic misfit η can than be calculated from the following formula. 58

67 1 da η = (4.6) a dx Si where a is the lattice parameter for aluminium, and da dx Si (4.7) relates to the change of lattice parameter caused by the addition of silicon in the case that Vegard s law is applied. Equation 4.8 calculates the lattice parameter for the strained aluminium lattice in case silicon is added. R = R R 1 1 R N Si strained latice Al (4.8) Al RAl and RSi are the lattice parameters for aluminium and silicon, N is the mole fraction of silicon. The mole fraction in front of the liquid according the right hand graph of figure 3.15 is The lattice parameters for silicon and aluminium are 5.43 and 4.5 Ǻ respectively. These figures will result in a misfit factor for this system of 0.6%. Putting all these data into formula 4.5 will result in the energy available as a driving force to move the liquid film in two crystallographic directions, resulting in: G 100 = -294 J/m 3 and G111 = -404 J/m 3. Although very small, the coherency energy is negative, meaning that it can provide in theory, the energy for the liquid film to move Chemical free energy The chemical free energy or Gibbs energy change in the system is caused by the change in concentration of the main three alloying elements namely; Mn, Cu and Si. The difference in free energy of the system in front of and behind the liquid film should provide an answer as to whether there is a sufficient surplus of energy to maintain the movement of the liquid film. The chemical free energy of the alloy at the two locations has been calculated by means of the software program JMatPro-4.0. The actual concentration of the copper containing core alloy with the measured concentration as taken from the left hand graph of figure 3.16 has been used as input. Table 4.2 shows the outcome of the free energy calculation. 59

68 Table 4.2: Element concentrations in wt% used for free energy change calculations. Si Mn Cu Fe Al G(J/mol) at 595 C Ahead of the liquid Balance film(core 2) Behind the liquid film (3.16) Balance Free energy change 560 According to thermodynamics the positive free energy change indicates that it will not be a spontaneous reaction and that it is unlikely that the chemical free energy change is the driving force behind liquid film migration during aluminium brazing Curvature driven (sub) grain coarsening As discussed in chapter 3, LFM seems to occur in pre-stretched materials that do not recrystallize at brazing temperatures. Previous studies [8,27] have revealed that samples stretched 10% prior to brazing had a recovered sub-grain structure after brazing with a subgrain size of around 1-2 µm. The area behind the liquid film is essentially dislocation and sub grain boundary free, since a liquid film that has passed through is not capable of supporting the existence of dislocations. The reduction in sub-grain or grain area is a possible source of energy for the movement of a liquid film. The amount of energy released by the elimination of sub-grain or grain boundaries is of course dependent on the size and orientation of the sub-grains or grains. Sub-grain boundaries and low angle boundaries can be represented by an array of dislocations [28]. The grain boundary energy between two (sub-)grains is dependent on the misorientation between the two (sub) grains. If the spacing of the dislocations of Burgers vector b in the boundary between two adjacent grains equals h, then the crystals on both sides of the boundary show a misorientation of θ. The relation between misorientation and grain boundary energy is represented by the Read and Shockley equation [28].: γgb = E0 θ( A0 lnθ) (4.9) where Gb E 0 = and 4π(1 υ) A 0 b = 1+ ln (4.10) 2πr 0 G is the shear modulus, υ is the Poisson ration, b is the Burgers vector and r 0 is the radius of the dislocation core. The energy between two grains is dependent on the misorientation between the two grains as illustrated by the following figure. 60

69 Energy Boundary energy Misorientation( ) Fig. 4.9: The energy of a grain boundary as a function of misorientation. Grain boundaries are defined as being either Low Angle or High Angle Grain Boundaries. High Angle Grain Boundaries (HAGB) are those exceeding a misorientation [29]. The HAGB has the highest boundary energy of J/m 2 for aluminium [28]. The total energy represented by a (sub-)grain surface is dependent on the size and the misorientation between the grains. In the calculation of the amount of surface energy available, the grain boundary energy for an HAGB has been used. Figure 4.10 represents the amount of grain boundary energy assuming cubic crystals. The total grain boundary surface energy is calculated from the formula (4.11) [28]: E gb = 3 γ D gb (4.11) where D is the (sub) grain diameter and γgb the grain boundary energy J/m (sub) grain bioundary energy (kj/m 2 ) (sub) grain diameter (µm) Fig 4.10: Total (sub) grain boundary energy. The amount of energy drops to low levels at larger (sub) grain sizes. If the (sub) grain size of 2 µm as reported by Yang [8] is used, then the amount of energy available would be -486 kj/m 3, however the true amount of energy is less since not all grain boundaries will be HAGB s. 61

70 4.5.4 Grain boundary wetting In a liquid-solid system wetting of the grain boundaries will occur if the following condition is fulfilled: γgb > 2γ sl (4.12) Wetting would occur if the grain boundary energy between two grains is larger than the sum of the two newly created solid-liquid interface energies. The thermodynamic explanation for wetting is explained by the figure γ 2γ sl γ gb T w T Fig Wetting of a grain boundary [29]. The grain boundary energy between two grains depends on the misorientation between the two grains [28]. In the brazing system studied, the grain boundary energy is γ gb =0.324 J/m 2, the surface tension at the interface between the liquid 10% silicon alloy and solid pure aluminium at 595 C would according to Gündüz and Hunt [30] be γ sl =0.169±0.20 J/m 2. Due to the uncertainty of the measurement, the value for γ sl ranges between wetting and nonwetting conditions. Assume that the average value for γsl is valid, then the total free energy released by wetting of the grain boundaries is a function of the total grain boundary surface, which obviously is related to the grain size. The same assumption can be used as in the previous calculation, that all grain boundaries are HGBA s. Figure 4.12 represents the energy available when a material with different grain sizes has been wetted with the 10% aluminium silicon alloy having a surface tension within a standard deviation of the measurement of Gündüz and Hunt [30]. 62

71 200 Energy (kj/m 3 ) Av.. ysl=0.169 γ sl = J J/m2 / m 2 Min.. ysl=0.149 γ = J J/m2 / m 2 Max.. γ ysl=0.189 = J J/m2 / m grain size (µm) Fig. 4.12: Energy release by grain boundary wetting. Depending on the value of the surface tension of the liquid wetting the grain boundaries, energy will be released or consumed Reduction of particle/matrix interface energy In the previous section the energy release because of (sub) grain boundary reduction was calculated. In the AA3XXX series core alloy under investigation, a second surface energy source is available namely the surface energy between precipitates and the core alloy matrix. According to Li and Arnberg [31] and Yamauchi and Kato [27] the amount and size of precipitates are dependent on the thermomechanical treatment of the alloys prior to brazing. For the calculation of the amount of particle-matrix interface energy, the precipitate size and volume fraction have been varied. The volume fraction of precipitates ranges between 0.2 and 1.0 volume percent while the size of the precipitates range between 20 and 200 nm diameter. The calculations are represented in figure 4.13 and for which it is assumed that the precipitates are mainly MnAl 6. The precipitate interface energy between the matrix and precipitate according to Chen [32] is given as 0.26 J/m 2. The formula used to calculate the interface energy is: E 6 = γmnal (4.13) 6 D interface F v D = precipitate diameter, F v = volume fraction, γ MnAl 6 = 0.26 J/m 2 63

72 800 Total interface energy (kj/m 2 ) Fv=0.2% F v Fv=0.4% F v Fv=0.6% F v Fv=0.8% F v Fv=1.0% F v precipitate diameter (nm) Fig. 4.13: Total interface energy between matrix and precipitates. With an average precipitate size of 100 nm, the total available interface energy ranges between and -156 kj/m 2 for precipitate volume fractions between 0.2 and 1.0% Surface tension anisotropy In 1998, Kirkaldy [33] discussed the validity of the coherency strain energy [5,6] as the driving force for LFM. The conclusion from the discussion was that it was not the coherency strain energy but a weak anisotropy at the two solid-liquid interfaces that was responsible for the interface mobility. Unfortunately, this conclusion was not corroborated by any observation and/or calculations. One can, however, argue that in the brazing system being studied, at the two interfaces, liquid film concentration differences of the alloying elements will occur. This is represented by figure d Affected Area Si 1.18 Mn 0.17 Cu 0.08 Fe 0.01 Si 10 Cu<0.4 liquid γ sl Original Core Si 0.09 Mn 1.01 Cu 0.4 Fe 0.2 Fig. 4.14: Concentration differences at the liquid interfaces. 64

73 Fe and Mn are forming precipitates in the liquid film while copper remains in the liquid. The liquid originating from the molten clad alloy does not contain any copper; hence the copper concentration in the liquid is the highest at the interface with the original core alloy and the lowest at the side where there is hardly any copper in solid solution. Elements going into solution will in most cases lower the solid-liquid surface tension of the bulk liquid [34]. However, due to the low concentration of copper the lowering effect would be minor, but might be sufficient to create a driving force. This argument is only valid for the experiments with the copper containing core alloy. This would mean that in order for surface tension anisotropy to play a role, an element other than copper would have to create a concentration difference between the two interfaces of the liquid film. It is therefore unlikely that the surface tension anisotropy plays a significant role in the onset of liquid film migration, since LFM does occur in both alloys studied in this thesis. 4.6 Discussion Liquid Film Migration From the diffusion measurements it seems most likely that the profiles measured can only originate from a moving boundary. Another explanation for the observed profiles is that (sub-) grain boundary diffusion plays a significant role. In the overview by Mishin et al. [35] the fundamentals of grain boundary diffusion were discussed. There are three types of grain boundary diffusion recognized that can result in three different profiles. In figure 4.15 the three types are schematically illustrated. Fig. 4.15: Schematic illustration of type A, B and C diffusion kinetics. 1 Type A is dominant when ( D eff)2 is greater than the spacing (d) between two boundaries. 65

74 1 ( D eff t) 2 >> d (4.14) D eff b ( f) Dv = f D + 1 (4.15) where f is the volume fraction of grain boundaries in a polycrystal, D b is the grain boundary diffusion coefficient and D v is the volume diffusion coefficient. Type A results in a uniform diffusion profile similar to that measured in this study ahead of the interface as represented in fig If for the sake of argument it is assumed that the measured profile is to originate from type A diffusion kinetics, it would be possible to calculate the distance between two boundaries. If next the assumption is that the sample consists of equiaxed grains then this distance would be the actual grain size. For the calculation, the effective diffusion distance is used as represented in fig. 4.3 where S eff equals 41 µm, S eff also satisfies the equation: S eff 1 ( Deff t)2 = (4.16) Equation 4.15 allows us to calculate D eff which is in this case will be 1.4x10-11 m 2 /s. This can be used as input for 4.14 where according to Mishin [35] for Face Centred Cubic (FCC) materials like aluminium f = D b is 4.14x10-10 m 2 /s and Dv is 1.16x10-12 m 2 /s, resulting in The volume fraction of grain boundaries is: δ f = (4.17) d Where δ is the grain boundary width, being 0.5 nm [32], resulting in a grain size d of 16 nm. This grain size for deformed aluminium is extremely small. Yang [8] reported 2 µm sub grains after brazing in deformed material which is orders of magnitude larger. Even a severe deformation process such as Equal Channel Angular Pressing could only achieve sub grain sizes of around 300 nm at applied strains between 8 and 13 [36] which is far more than the strain ranges applied to the measured samples. Another argument in favour of a liquid film changing the interface between clad and core alloy during brazing is the presence of particles and precipitates at the interface. It would be challenging to find a mechanism and driving force capable of moving solid particles through a solid matrix. 66

75 4.6.2 Driving forces A number of observed changes could provide the necessary energy for a mechanism to occur. Table 4.3 summaries the origin and size of the energies available in the system: Table 4.3: Potential available energy sources Nature G (J/m 3 ) Coherency strain energy -300,-400 Chemical energy 55.9 M (Sub) grain coarsening (d=2µm) -486 k Grain boundary wetting (d=2µm) -100 k Particle/matrix interface (F v 1% d=100nm) -156 k Surface tension anisotropy -<1k If the energy from coherency strain would play a role, all samples would suffer from LFM. In all samples diffusion of silicon in front of the interface takes place; which according Yoon [6] should be the driving force for LFM. This is however not observed in most samples of this study; therefore the coherency strain is considered not to play a significant role in the onset of LFM. Table 4.3 shows that the chemical changes cannot provide the driving force, as the changes would consume energy. Generally, in all samples given a specific homogenisation treatment of both core alloys, the size and number of precipitates will not change. Therefore, the available particle/matrix energy would be the same for all these samples. The size and number of precipitates can play a role in stabilising a (sub-) grain structures in the strained samples at elevated temperature retaining more energy available for the onset of LFM. This will be discussed in a chapter 5. Of the other three sources listed, the most likely source to play a major roll in the onset of liquid film migration is (sub-)grain coarsening or the energy associated with dislocations. In chapter 3 none of the samples from both core alloys irrespective of the homogenisation treatment, shows any signs of LFM when brazed in the full O temper (soft) condition. At increasing applied strains, some of the materials start to show a metallurgical interaction at the interface between the core and clad alloy as described in chapter 3. The interaction consumes part of the core alloy as can be seen in figures 3.11 and The interaction will stop once the grain structure of the core alloy changes; which is at the onset of recrystallization. However, the interaction between the core and clad only occurs within in a certain window of applied strains where recrystallization of the core alloy does not occur. The same observation was made previously by others in the past [8,20,27]. The two metallurgical processes LFM and recrystallization are in competition with each other. 67

76 Available dislocation structures or sub-grains seem to play an important role in the onset of LFM. The same observation has been made for systems other than aluminium brazing [37,38]. Whether it is the energy released by eliminating the dislocation structure itself or the penetration by wetting of the molten clad alloy along this sub-structure still needs to be clarified. Sub-grain boundaries are considered to have a low angle misorientation [28], the energy associated with the boundaries might not be enough to fulfil equation 4.12 allowing wetting, where in that case the energy stored in the sub-structure will be consumed by a different mechanism allowing the liquid film to move. 4.7 Conclusion There is ample evidence that the changes occurring during brazing are caused by a liquid film moving from the interface between the molten clad and the core alloy into the core alloy. The key for the onset of LFM is the presence of dislocations, whether or not ordered in a sub-grain structure. Strained samples that have not recrystallized on reaching the brazing temperature are the most susceptible to LFM. In the next chapter the actual role of these dislocations and a possible mechanism behind LFM will be investigated and discussed. 4.8 References 1. D.N. Yoon and W.J. Huppmann, Acta Metall. Vol. 27, 1979, pp J.J. Pepe and W.F. Savage, Weld. Res. Sup., December 1970, 545s-553s. 3. W.J. Huppmann, Z. Metallkde. Bd. 70 (1979) H. 12, pp M. Hillert and G. R. Purdy, Acta Metall. Vol , pp D.N. Yoon, Annu. Rev. Mater. Sci D.N. Yoon, Int. Mater. Rev., 1995, Vol.40 No. 4, pp S. Annavarapu and R.D. Doherty, Acta. Metal. Mater. Vol 43. No.8, pp , H.S. Yang, and R.A. Woods, SAE paper V.L. Acof and R.G. Thompson, IBM J. Res. Develop. Vol 44, No 5 September G.B. Schaffer,S.H. Huo, J. Drennan and G.J. Auchterlonie, Acta Mater. 49 (2001) E.D. Manson et al. Acta Mater. 50 (2002) W.A.G. McPee, G.B. Schaffer and J. Drennan, Acta Mater. 51 (2003) T. Antonsson, et all, Int. J. of Refr. Met. & Hard Mat. 21 (2003) O. Schenk and J.L Uria, J. Metamorphic Geol, 2005, 23, U. Chntababu, V.R. Chary and S.P. Gupta, Can. Metall. Quart., vol.46, No. 2, pp , S.W. Barker and G.R. Purdy, Acta Mater. Vol. 46, No. 2, pp ,

77 17. Randall M. German, Liquid Phase Sintering, 1985 Plenum Press, New York, ISBN O.A. Ojo, N.L. Richards, M.C. Chaturvedi, Scr. Mater. 50 (2004) C. Huang, S Kou, Supp. to the Weld. J., April 2004, 111-s, 122-s. 20. A. Fukumoto, T. Doko, Proceedings of the 9 th International Conference on Aluminium Alloys (2004), J.S. Ryu, M. S. Kim, D. Jung, J. of Mat. Pro. Technol (2002) W.A. Tiller, K. A. Jackson, J.W. Rutter, and B. Chalmers, Acta Met., vol. 1, July 1953, Y.Du et all, Mater. Sci. and Eng. A ) M.S. Sulonen, Ann. Acad. Sci. Fenn. A, No.4 (1957), M. Hillert, G. R. Purdy, Acta Metall., Vol. 26 pp P. M. Sutton, Physical review, vol. 91, No. 1, August 13, 1953, pp S. Yamauchi, k. Kato, Keinkizoku 1991, vol. 41, no.4, pp F. J. Humphreys and M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon, Oxford, 1996, ISBN R.D. Doherty, D.A. Hughes, F.J. Humprhreys et al, Mater. Sci. and Eng. A238 (1997) M. Gündüz, J.D. Hunt, Acta Metall. Vol. 33, No. 9, pp , Y.J. Li and L. Arnberg, Acta Mater. 51(2003) pp S. Chen, Thesis, ISBN , J.S. Kirkaldy, Acta. Mater. Vol. 46. No. 14 pp , A.R. Miedema, F.J.A. den Broeder, Z. Metallkde, Bd. 70 (1979), pp Y.Mishin, Chr. Herzig, J. Bernardini and W. Gust, Inter. Mater. Rev. 1997, vol.42, no. 4, R. Kapoor, J.K. Chakravartty, Acta Mater. 55 (2007) W. Schatt, W.A. Kaysserr, S. Rolle, A. Sibillia, E. Friedrich and G. Petzow, Pow. Metall. Int. vol.19, no.1, 1987, H. Fredrikson, A. Eliasson and L. Ekbom, Int. J. of Refr. Met. & Hard Mater. 13 (1995)

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79 Chapter 5 Strain Induced Liquid Film Migration 5.1 Introduction From the previous chapters it became clear that the occurrence of Liquid Film Migration is linked to the occurrence of recrystallization of the core alloy. When the core alloy recrystallizes before the melting temperature of the clad alloy is reached, no LFM is observed. At low levels of deformation as in the case of the stretched samples ε <0.1, the dominant recrystallization mechanism in aluminium is Strain Induced Boundary Migration [1]. SIBM involves the bulging of a pre-existing grain boundary, leaving a dislocation-free region behind the migrating boundary. The driving force for SIBM is the difference in dislocation density between opposite sides of the grain boundary. SIBM can lead to preferential growth of a few grains resulting in material with extremely large grains. These large grains have all been observed in the samples in the figures After recrystallization all grains in the core alloy are large and extend across the full thickness of the core. In the non-recrystallized samples LFM occurs at certain strain levels. The purpose of this chapter is to prove that at the interface of the clad alloy and core alloy a similar process to SIBM can take place, resulting in the movement of a liquid film driven by the difference in dislocation density between both sides of this liquid film. The process in which a liquid film is caused to move because of these differences is called Strain Induced Liquid Film Migration. During brazing SIBM and SILFM are both in competition with each other. 71

80 5.2 Strain Induced Grain Boundary Migration SIBM is generally to be held responsible for the change in grain structure at low strain levels. Bailey and Hirsch [2] were the first to analyze the kinetics of the process. At a later stage, Bate and Hutchinson [3] re-evaluated the proposed kinetics and suggested some correction to compensate for a geometrical effect. However, the basis of the kinetics remained the same. According to Bailey and Hirsch [2] an original boundary is free to bulge under the action of a differential driving force P= P 1 P2 where P 1 and P 2 are the stored energies of the two adjacent grains. For nucleation to occur: R0 2γ gb/ P (5.1) or where γ gb is the specific energy of the grain boundary and assuming P to be constant, R 0 2γ = P(1 gb f) (5.2) f is the difference in dislocation density on opposite sides of the grain boundary. If the dislocation densities differ by 10%, then f=0.9. However, due to the bulging, the pressure on the concave side ( P 2 ) continuously decreases while the pressure on the convex ( P 1 ) side remains constant. Therefore P is not constant. Figure 5.1 shows how SIBM evolves [3,4]. Fig 5.1: a) SIBM of a boundary separating a grain with a low pressure( P 2 ) from one with a higher pressure ( P 1 ). b) Dragging of the dislocation structure behind the migrating boundary. c) The migrating boundary is free from dislocations. The correction made by Bate and Hutchinson [3] for the non-constant pressure over the bulging grain boundary resulted in the following equation: 72

81 R 0 = P 2γ gb (1 f) (5.3) Equation 5.3 relates to a critical radius that will bulge into a recrystallization nucleus to the grain boundary energy and the difference in dislocation density on opposite sides of the critical area of the bulge. Equation 5.3 is used to explain that a liquid film wetting a grain boundary can cause a nucleus to grow at lower dislocation densities compared to a nonwetted grain boundary. 5.3 Strain Induced Liquid Film Migration Assume a liquid metal has wetted a grain boundary. How and if a grain boundary is wetted will be discussed later. The rigid structure of the solid grain boundary has now been displaced by a liquid layer. Hence the interface energy between the two grains is no longer represented by grain boundary energy ( γ gb ), but by the energy of the solid liquid interface ( γ sl ), which has a much lower value. The High Angle Grain Boundary (HAGB) energy in aluminium is J/m 2 [4] while the Solid-Liquid interface energy between solid aluminium and a liquid AlSi 10 wt% alloy is J/m 2 [5]. In figure 5.2 the effect of the interface energy is shown, when equation 5.3 is used to calculate the critical area to bulge for different dislocation densities or energy levels in P Critical radius R γgb γ f f=0.9 γsl γ f f=0.9 γsl γ f f=0.95 γsl γ f f=0.98 γsl f f=0.99 γ sl Presure grain P 1 Fig. 5.2: Effect of interface energy on critical radius to bulge. In fig. 5.2 shows when the surface energy is lowered as in the case if γgb goes to γ sl and all other parameters stay the same, there is a reduction in the critical radius ( R 0 ) (vertical line in 73

82 fig. 5.2). Also it can be seen that the pressure (energy) in the grain with the highest dislocation density ( P 1 ) can be lower in the case of γ sl and still result in the same critical area as with the higher energy in grain ( P 1 ) in combination with γ gb (horizontal line in fig. 5.2). Equation 5.3 show that there is a linear relationship between the interfacial energy and the energy available in grain P 1 to cause boundary movement in the situation where f and R 0 remain constant. In all cases a lower interfacial energy between grains would result in a movement of the boundary at a lower stored energy in the grains and also with a smaller difference in energy between the two grains. In the alloy systems studied in this thesis, a competition between SILFM and SIBM can take place. Before SILFM can take place, a liquid phase has to form and penetrate the grain boundaries. This liquid phase comes from the AA4XXX series clad layer on top of the core alloy. The clad alloy starts to melt at 577 C. When the core grain boundaries have been penetrated and sufficient energy is still stored in the grains, SILFM can occur. Below 577 C when no liquid phase is present only SIBM can occur. All observations in chapter 3 show that the interaction between the molten clad alloy and the core alloy, now called SILFM, only occurs when the core does not show signs of recrystallization by SIBM. SIBM takes away the energy required to start SILFM. In a number of samples neither SILFM nor SIBM is observed, especially in the samples with low levels of stretching up to 3%. At these levels the energy present in the system is too low to cause either SILFM or SIBM. As has been presented in chapter 3, the composition of the area behind the migrated film is at the thermodynamic equilibrium corresponding to the temperature of the experiment [6]. The silicon concentration in the solid behind the liquid film is 1.2 wt% compared to 0.09 wt% in front of the liquid film. The liquid film itself will, according to the phase diagram, contain approximately 12 wt% Si. The increase in silicon in solid solution is coming from the liquid film, which if it wants to stay liquid has to be replenished with silicon from the molten clad layer. The concentration difference between the front of the liquid film and the film close to the clad layer assures diffusion of silicon to the silicon depleted film, and possibly even a liquid transport to balance the silicon difference. Figure 5.3 shows the silicon distribution map of the liquid film both in front and in the wake of the liquid film. 74

83 Silicon feed from clad alloy Wake Si Liquid film Front 50 µm 50 µm Fig. 5.3: Silicon distribution. The changes occurring due to SILFM are identical to the normal LFM as described by Yoon [7]. However, the driving force is not caused by the coherency strain generated by the diffusion of elements in front of the liquid film, but by the difference in strain between the new grains and those already present due to the stretching prior to brazing. 5.4 Grain boundary wetting A prerequisite for SILFM to occur is the wetting of the grain boundaries by the molten clad alloy. As stated earlier in chapter 4, wetting occurs when the sum of energy of the two wetted newly created surfaces is smaller than the energy of the original non wetted grain boundary. The grain boundary energy depends on the misorientation (θ ) between the two adjacent grains [8] and is given by: ( lnθ) γgb = γ0 θ A (5.4) where γ 0 = Gb 4π(1 υ) and b A= 1+ ln 2π r 0, G is the shear modulus, b is the Burgers vector, υ is the Poisson ratio and r 0 is the radius of the dislocation core usually between 1 and 5 times the Burgers vector. Equation 5.4 can be used to calculate at which misorientation equation 4.12 is fulfilled for a certain liquid-solid surface tension. It can be rewritten to: Gb e γgb = θ ln (5.5) 4π(1 υ) 2πθ 75

84 According to Humphreys and Hatherly [4] equation 5.5 is valid at misorientations up to 15 and can be re-written to: γ θ θ ( HAGB) 1 ln (5.6) θm θ gb = γgb m Where γgb(hagb) is J/m 2 and θ m =15. Equation 5.6 is represented by figure grain boundary energy (J/m 2 ) mis-orientation ( ) Fig. 5.4: Relationship between grain boundary energy and grain misorientation. To fulfil the condition of equation 4.12, only grain boundaries with a misorientation of 15 or higher will be wetted by a liquid aluminium-silicon alloy as measured by Gündüz and Hunt [5]. Therefore SILFM will start at wetted high angle grain boundaries. Low Angle Grain Boundaries (LAGB s) which originated from the dislocation network present in a deformed aluminium alloy, as in our case, cannot be wetted. Although the liquid aluminium-silicon alloy can only wet an HAGB it does not automatically mean that a possible network of dislocations (sub-grain structure) is not present in the deformed structure. It only means that this substructure cannot be wetted by the liquid clad alloy. This will restrict the number of possible mechanisms responsible for the observed changes. 5.5 Experimental setup The experiments were first set up to prove whether grain boundary wetting can occur in the studied clad-core system and secondly if the observed changes can be caused by SILFM. The samples studied were either furnace brazed or salt bath treated. 76

85 All cross-sections were made perpendicular to the rolling and stretching direction Grain Boundary Wetting To determine if the molten clad alloy could wet the core alloy, brazed samples were first prepared and polished according to standard procedures for aluminium. Subsequently the samples were electrochemically etched by the following procedure to remove any residual strain in the surface: 6.4 g of potassium iodide was dissolved in 320 ml methanol. A 99.99% pure aluminium sheet was used as the cathode. The etched area of the sample was approx. 1 mm 2. The sample was the anode and was immersed in the methanol solution which was stirred by a magnetic stirrer. A potential of 7 volts DC and a current of between 0.01 and A was applied for 150 seconds. After etching the sample was thoroughly rinsed with demineralised water and ultrasonically rinsed in ethanol. After drying with warm air the sample was ready. An additional grain boundary investigation was carried out by using Gallium Enhanced Microscopy [9] (GEM). This technique should be able to reveal low and high angle grain boundaries by the progressive penetration of the liquid gallium along these grain boundaries. The high atomic mass of gallium (69.72) compared to that of aluminium (29.98) makes gallium an attractive element to study by SEM in the back scattering mode. Gallium has a strong preference for penetrating grain boundaries in aluminium polycrystals [10]. Two Scanning Electron Microscopes were used in this study. The initial experiments with gallium were conducted using a LEO 438vp SEM with EDAX TSL. The following gallium experiments were carried out on a ZEISS Ultra 55 Gemini FE-SEM with a Pegasus XM 4 Hikari system. This microscope was also used to investigate the silicon penetration from the clad alloy into the core alloy Strain Induced Liquid Film Migration The focus of the experiments in this section is to demonstrate the existence of energy in the system in the form of sub-grain boundaries or dislocations at the brazing temperature to facilitate SILFM. Some energy in the form of dislocations or sub-grains must be available in the core alloy structure to provide the necessary driving force for SILFM. Scanning Electron Microscopy and high resolution Electron Back Scatter Diffraction (EBSD) were applied to a number of samples. EBSD should provide detailed information on the orientation of the grains and possibly sub-grains. 77

86 The two alloys as presented in chapter 3 were used, which were given the cycle 2 homogenization, namely 24 hrs at 500 C. For the EBSD measurements, the following settings were used: - Aperture 120, high current - EHT = 20 kv - WD = 16 mm fps - Step size 0.55µm - Magnification 250x (field of view µm) - All samples were corrected for a 70 tilt Of the measured samples the Image Quality, the Confidence Index, the Inverse Pole Figure and the Kernel average misorientation were determined. The image quality describes the quality of the electron back scatter pattern. The Confidence Index (CI) quantifies the reliability of the indexed patterns; the CI ranges between 0 and 1. The Inverse Pole Figure displays the orientation of the different grains within the sample. Across the grains and within one grain the misorientation was measured to determine the possible existence of sub-grain boundaries. In order to evaluate the stored energy for a given point, the most appropriate quantity is the Kernel Average Misorientation (KAM). The KAM is defined for a given point as the average misorientation of that point with all of its neighbours [11]. KAM has been used by Takama et al. [12] to calculate the stored energy, which can be used to determine the possible onset of SIBM or SILFM. The GEM sample treatment was carried out as described by Hagström et al. [9]. 78

87 5.6 Results Grain boundary wetting by the silicon-rich cladding alloy residual clad alloy Fig 5.6a: Silicon penetration during brazing along grain boundaries of the [Mn] core alloy homogenised at 500 C for 24 hours and stretched 5% before brazing. The silicon-aluminium clad alloy seems to penetrate for at least 50 µm along the grain boundaries of the core alloy as can be seen in figure 5.6a. Figure 5.6b shows penetration of silicon from the cladding alloy into the [Mn+Cu] core alloy, which was homogenised for 24 hours at 500 C and stretched 5% prior to brazing. 79

88 residual clad alloy Fig 5.6b: Silicon penetration during brazing along grain boundaries of the [Mn+Cu] core alloy homogenised at 500 C for 24 hours and stretched 5% before brazing. Both samples show a strong penetration of silicon from the molten clad alloy into the core alloy along the grain boundaries. Further analysis at grain boundaries in front of the liquid film revealed very thin layers of silicon embedded between two grains. The etching applied to these samples only dissolves the matrix aluminium and elements in solid solution. However, the silicon film is so thin that even the smallest disruption in the etching bath will cause the silicon film to break Grain boundary wetting by Gallium Penetration of grain boundaries by gallium can be watched in situ, however the specimen requires very good preparation. The as-polished surface will result in a blurred image of the grain boundary, and some residual scratches of the polishing action are also still visible. Figure 5.7a shows the penetration of gallium along the grain boundaries but at the same time the already penetrated grain boundaries become blurred. 80

89 Fig. 5.7a: Progressive penetration of gallium. If however the electrochemical etching action is too strong, the surface becomes too rough and the gallium is hardly visible in the grain boundary. Figure 5.7b shows the interface between a just-polished area and electrochemically etched area. The penetration of the gallium is difficult to see because of the surface roughness caused by the electrochemical etching. Fig. 5.7b: Gallium penetration in mechanical and electrochemically etched material. To obtain useful results with GEM, special care is required when applying the electrochemical etching procedure. In case the procedure has been carried out correctly, detailed results can be obtained as presented in the figure 5.8. Fig. 5.8: Resolution of GEM, Mn-Cu alloy non-homogenised, stretched 5% prior to brazing. The resolution of the first set of results as presented in fig. 5.8, however, did not indicate the existence of a sub-grain structure as reported by Yang and Woods [13]. As presented in chapter 3, the Mn-Cu non-homogenised core alloy showed the strongest LFM and the core alloy does not seem to recrystallize even after stretching by 7 or 10%. It is most likely that if a sub-grain structure would have been developed it would be in these samples. 81

90 Interface Original core alloy Fig. 5.9: GEM, Mn-Cu alloy non-homogenised, stretched 10% prior brazing. Use of a high resolution SEM made a more detailed study possible but even this did not conclusively show a sub-grain structure. Figure 5.9 shows the gallium penetration from the interface between clad and core alloy into the core alloy. Figure 5.10 shows the details of the gallium penetration at a triple point. Sub grain boundary? Fig. 5.10: GEM high resolution. An interesting observation made by high resolution GEM is the apparent presence of a grain boundary in particles which can be seen in the right hand picture of fig Figure 5.11 shows three particles in which the grain boundary is clearly visible. A possible explanation is that the SEM images were taken in the Back Scatter mode, meaning that the high atomic weight of gallium will produce a very strong signal and since the particles are quite thin they might appear to be transparent to the underlying gallium filled grain boundary. Figure 5.11 shows another example of a grain boundary showing up under a particle. 82

91 Grain boundary Fig. 5.11: Grain boundaries under particles Strain Induced Liquid Film Migration SEM EBSD images were taken to try and obtain conclusive information on the presence of sub-grains in the deformed materials after brazing. Figures show the Image Quality, the Confidence Index and Inverse Pole Figure of alloy 80BC (Mn core alloy, homogenised for 24 hrs at 500 C) and alloy 83BC (Mn/Cu core, homogenised for 24 hrs at 500 C). Moreover, the misorientation profiles across grains and within a single grain are presented. To determine any residual strain within grains the Kernel Average Misorientation was measured. KAM represents the average misorientation of a given point with respect to its neighbours. The intensity is an indication of the difference in misorientation and shows that the grains are not completely strain free. 83

92 Fig. 5.12: Image Quality of post-braze 80BC. Fig. 5.13: Confidence Index of post-braze 80BC. 84

93 Fig. 5.14: IPF of pos- braze 80BC. Misorientation BC_0 80BC_1 80BC_3 Misorientation Misorientation Distance (µm) Distance (µm) Distance (µm) Misorientation BC_5 80BC_7 80BC_10 Misorientation Misorientation Distance (µm) Distance (µm) Distance (µm) Fig. 5.15: Horizontal misorientation profiles of 80BC from fig Red line is the point to point difference. Blue line is the point to origin difference. 85

94 Misorientation BC_0 80BC_1 80BC_3 Misorientation Misorientation Distance (µm) Distance (µm) Distance (µm) Misorientation BC_5 80BC_7 80BC_10 Misorientation Misorientation Distance (µm) Distance (µm) Distance (µm) Fig Vertical misorientation profiles of 80BC from fig Red line is the point to point difference. Blue line is the point to origin difference. Fig Kernel Average Misorientation of 80BC. 86

95 Fig. 5.18: Image Quality of post-braze 83BC. Fig. 5.19: Confidence Index of post-braze 83BC. 87

96 Fig. 5.20: IPF of post-braze 83BC. Misorientation BC_0 Misorientation BC_1 Misorientation BC_ Distance (µm) Distance (µm) Distance (µm) Misorientation BC_5 Misorientation BC_7 Misorientation BC_ Distance (µm) Distance (µm) Distance (µm) Fig. 5.21: Horizontal misorientation profiles of 80BC from fig Red line is the point to point difference. Blue line is the point to origin difference. 88

97 3 83BC_0 3 83BC_1 3 83BC_3 Misorientation 2 1 Misorientation 2 1 Misorientation Distance (µm) Distance (µm) Distance (µm) Misorientation BC_5 Misorientation BC_7 Misorientation BC_ Distance (µm) Distance (µm) Distance (µm) Fig. 5.22: Vertical misorientation profiles of 83BC from fig Red line is the point to point difference. Blue line is the point to origin difference. Fig. 5.23: Kernel Average Misorientation of 83BC. 89

98 5.7 Discussion Grain boundary penetration by silicon from the molten clad alloy The observations as presented in figures 5.6a+b make it clear that molten clad alloy rich in silicon has penetrated the grain boundaries of the core alloy. The grain boundaries in figure 5.24 have been highlighted and show that the silicon has travelled over a distance of over 75 µm. The penetration of the silicon from the molten clad alloy is difficult to observe in polished specimens due to the narrowness and fragility of the silicon film. Only in the electrochemically polished samples if a residue of the film remains sticking out of the grain boundary between two grains, it can be observed. Fig. 5.24: Depth of Silicon penetration. No previous publications are known describing or showing the penetration of silicon along the grain boundaries to such a depth as observed in this study. Grain boundary wetting occurs when the following condition of 2 γ is satisfied. sl γgb According to equation 5.6 only wetting of grain boundaries with a mis-orientation of 15 or higher will occur. This means that the silicon-rich liquid from the cladding alloy only penetrates along high angle grain boundaries. This liquid will not penetrate any sub-grain boundary since they have a low misorientation (<15 ). Therefore, studying the silicon penetration will not reveal the possible existence of a sub-grain structure. The observed penetration by silicon from the molten clad alloy closely resembles the characteristics of the intergranular film (morphology, thickness) which have been observed in a number of other liquid solid metal systems [14,15,16]. The following schematic structure 90

99 applies to the reported penetrations: first a groove is present with a length of the order of 10 µm, followed by a micrometric film which can be in the excess of several tens of microns and finally extending ahead of this is a nanometric film. Figure 5.25 shows the three different features of the silicon penetration. Fig. 5.25: Characteristics of grain boundary penetration Grain boundary penetration by Gallium In the observations made by GEM there seems to be no indication of the existence of subgrain boundaries, or at least in the samples studied. According to Hagström et all [9], GEM is capable of detecting sub-grain boundaries with a misorientation down to less than 1. In the case of this small misorientation it was mentioned that decoration of low angle boundaries does not take place anymore by ordinary wetting but through dislocation pipe diffusion. This type of diffusion is much slower than the wetting process but when the sub-grain boundaries are closely spaced the diffusion distances to be covered are within the time of the observation. The sub-grain size is dependent on the type of alloy and processing. In the study by Hagström, the average sub-grain size was 1-2 µm after 50 % cold rolling and annealing at 250 C for 32 hours. In our case the maximum equivalent reduction was 10 % and annealing took place at 595 C for several minutes. This practice would result in much larger sub-grains if the data of Sanström et al. [17] were to be extrapolated, with the result that decoration by pipe diffusion of gallium along sub-grain boundaries would be sporadic. In figure 5.10 a possible sub-grain boundary is just barely visible. One major drawback of GEM is the destructiveness of the gallium penetration. With prolonged exposure, liquid metal embrittlement can occur as illustrated by the figure

100 Fig. 5.26: Liquid metal embrittlement. In general, GEM can be used for studying grain boundaries and sub-grain boundaries. However, if no sub-grain boundaries are present or sufficiently developed in the samples, GEM does not provide any additional information compared to standard SEM Strain Induced Liquid Film Migration From all the data measured by EBSD no conclusive evidence for the existence of sub-grains could be found. The good image quality and high confidence index should allow the detection of sub-grains. However, a closer look at the IPF images of figure 5.14 sample 80BC_5 showed that a number of grains show some internal structure, not well organised but an indication of possible residual strain. Figure 5.27 shows an enlargement of fig BC_5 in which a faint indication of an organised sub-structure is visible. Fig. 5.27: Sub-structure within grains of 80BC_5 (TD= Tensile Direction). 92

101 A similar but better developed sub-structure was also observed by Hurley [18] in a 20 % cold rolled Al-0.13 wt% Mg alloy. The material studied had an equivalent cold rolling of 10 % and was annealed at 595 C. Sub-grain boundaries are considered to have a misorientation of 0.5 or over. No clear subgrain boundaries could be detected. Figure 5.28 shows the misorientation between the steps of 0.55 µm within a grain of sample 80BC_5. There are misorientations present, more than in the grains of 80BC_0 or 80BC_10, but not enough to be clearly identified as sub-grains. Fig. 5.28: Sub-structure within a grain of 80BC_5. The KAM measurements as presented in figures 5.17 and 5.23 are an indication of a disturbed internal grain structure. There seems to be a correlation between recrystallization and intensity of the KAM measurement. Figure 5.29 shows the grain structure and KAM of both alloys just before and after recrystallization. The recrystallized samples show a much lower intensity, indicating the removal of strain from the structure. 93

102 Fig. 5.29: KAM & Confidence index measurements on non- and recrystallized samples. Recrystallization caused by SIBM will reduce the internal strain within the grains. When the normal high angle grain boundary is replaced by a nanometric liquid film from the clad alloy, a reorganising process similar to SIBM can take place. The results presented show that indeed the onset of Liquid Film Migration is linked to SIBM. Figure 5.30 shows schematically how SIBM and SILFM are linked. In the graph the only difference between the two processes is the surface energy. P is the difference in energy between two grains. A slightly deformed material will recover when heated up. Recovery will reduce the amount of dislocations within the grains therefore recovery will reduce the energy available to cause SIBM. SIBM will also reduce the amount of energy available to cause SILFM, In principal recovery, SIBM and SILFM are all using the same source of energy present as dislocations available in the grains. Recovery is a process that can be described as a reorganisation and reduction of dislocations during a heat treatment. Anything that inhibits dislocation motion or annihilation will affect the recovery kinetics. To which extent recovery takes place in an alloy depends on time and temperature. Dislocations rearrange faster at higher temperatures. Strained aluminium brazing sheet can follow one of the three paths as presented in figure Material 1 strained to a certain level will recover until it reaches the conditions to fulfil equation 5.3. The material will now recrystallize by SIBM. Material 2, is less strained and will also recover but nothing will happen until the melting temperature of the clad alloy is reached. Due to the much lower surface tension of the liquid infiltrated grain boundaries, movement similar to material 1 will take place, the only difference being that less energy is required. The energy of the system 94

103 is now lowered by Strain Induced Liquid Film Migration. Material 3 is little strained and does not contain sufficient energy to have either SIBM or SILFM initiated. Fig. 5.30: SIBM and SILFM in relation to the temperature. 5.8 Conclusions Recrystallization in slightly strained materials takes place by Strain Induced Boundary Migration. In this process no new grains are formed, but some existing grains will grow at the expense of others. The main driving force is the total energy in the system and the difference in energy between two adjacent grains. The energy present in the system studied is a residue of the straining prior to brazing. This residual energy is the driving force for both SIBM and SILFM. For SILFM to occur, the liquefied clad alloy has to penetrate the grain boundaries of the core alloy. The liquid will only penetrate grain boundaries with a misorientation of 15 or over. Gallium is supposed to wet grain boundaries with a misorientation down to 0.5. However, low angle grain boundaries were not conclusively observed and neither did EBSD measurements reveal any well developed sub-grain structure. KAM measurements showed that there is still some energy left in the system although not in an organised structure. This energy at some levels is sufficient for SIBM or SILFM to occur. The critical parameter in SIBM is the high angle grain boundary energy. If this parameter is replaced by the lower value of the surface tension of the solid-liquid interface between the liquid clad alloy and solid grain, movement of this liquid filled grain boundary can take place at lower energy levels compared to the solid grain boundary. 95

104 This movement of a liquid boundary will cause compositional changes in the core alloy and will result in a compressed diffusion profile ahead of the liquid film. In the next chapter a more theoretical approach will be made to explain the compressed diffusion profile based on available literature. Al possible relationships between the occurrence of SILFM and thermo-mechanical processing of the core alloys will be discussed as well. 5.9 References 1. P.A. Beck and P.R. Sperry, J. Appl. Phys. 21, 150 (1950). 2. J.E.Bailey and P.B. Hirsch, Proc. R. Soc. Lond. A267, 11, (1962). 3. P. Bate and B. Hutchinson, Scr. Mater. 36, (1997) F. J. Humphreys and M. Hatherly, Recrystallization and Related Annealing Phenomena, Pergamon, Oxford, 1996, ISBN M. Gündüz, J.D. Hunt, Acta Metall. Vol. 33, No. 9, pp , S.W. Barker and G.R. Purdy, Acta Mater. 46, (1998), D.N. Yoon, Annu. Rev. Mater. Sci W.T. Read and W. Shockly, 1950, Phys. Rev. 78, J.Hagström, O.V. Mishin and B. Hutchinson, Scr. Mater. 49 (2003) E.Pereiro-Lopez, W. Ludwig and D. Bellet, Acta Mater., 52 (2004) S. Wright, D. Field and D. Dingley, Electron Backscatter Diffraction in Materials Science, Kluwer Academic, New York, (2000),ISBN X. 12. Y. Takayama, J. Szpunar, H. Kato, Materials Science Forum Vols (2005) pp H. S. Yang and R. A. Woods, SAE paper B. Straumal, T. Muschik, W. Gust, P. Predel, Acta Metall. Mater. 40, (1992), N. Marié, K. Wolski, M. Biscondi, Scri. Mater. 43 (2000) D. Chatain, E. Rabkin, J. Derenne, J. Berdardini, Acta Mater. 49 (2001) R. Sandström, B. Lehtinen, E. Hedman, I. Groza, S. Karlsson, J. Mater. Sci. 13 (1978) P. Hurley, F. Humphreys, Acta Mater. 51 (2003)

105 Chapter 6 Theoretical considerations of SILFM 6.1 Introduction In the previous chapter it has been demonstrated that the observations made can be related to a modification of the well known recrystallization process Strain Induced Boundary Migration. SIBM can be considered as a recrystallization process without the need for the formation of nuclei to start recrystallization [1]. Replacing a High Angle Grain Boundary by a liquid film results in the movement of this liquid film at lower strain levels than is the case for a normal solid grain boundary. This liquid film performs the same task as a normal HAGB, it removes dislocations, but due to its liquid state, it will alter the chemical composition of the swept area, bringing the chemical composition of the swept area into equilibrium at the temperature of the experiment. The driving force for the liquid film to move is generated from the excess energy stored in the strained grains. The previous chapters were dedicated to describing and evaluating the experimental data in a qualitative way. In this chapter an attempt will be made to quantify the observations. During Strain Induced Liquid Film Migration (SILFM) several processes take place simultaneously, i.e.: 1. Migration of a liquid film 2. Dissolution in front of the liquid film 3. Re-solidification behind the liquid film 4. Transportation of material from the front to the back of the liquid film 5. Transportation of liquid from the residual clad alloy to the liquid film 6. Diffusion from the liquid film into the core alloy 7. Diffusion from the core alloy to the liquid film 97

106 Figure 6.1 shows schematically the different processes occurring during SILFM. Fig 6.1: Different processes during SILFM. In this chapter a model will be proposed that fits with the observations made in this study. The basis of the model is the relation presented by Annavarapu and Doherty [2] to describe liquid phase sintering of spray cast materials. Their basic model includes most of the processes in figure 6.1, the only one not incorporated in their model is the transportation of liquid from the molten clad to the front of the liquid film. The driving force for the movement of the film is the reduction in grain surface energy. No reference is made to a possible second source of energy like strain energy. The model proposed will be discussed in view of the observations made for this thesis. Alternative models will also be addressed. 6.2 A simple model to describe SILFM After recrystallization has started, the high angle grain boundary will start to migrate. The driving force for grain boundary migration is represented by the following equation 2 α( ρf ρb) Gb 2γb 3Fvγb P= ( PDf PDb) + PC Pz = + (6.1) 2 R d 98

107 ( PDf PDb) is the driving pressure for growth where α is a constant of 0.5, ( ρf ρb) is the difference in dislocation density in front of and behind the high angle grain boundary, G is the shear modulus and b is the Burgers vector. P C is a result of the growth of a grain with radius Rinto the deformed structure and is given by the Gibbs-Thomson relationship. The total grain surface area is reduced during grain growth. P Z is a retarding pressure resulting from Zener pinning where particles with a diameter d and γb is the high angle grain boundary energy. F v is the volume fraction of As demonstrated in chapter 4, the area behind the liquid film has been re-deposited from the liquid phase and a liquid phase does not support the existence of dislocations. Therefore, the area behind the liquid film is dislocation free and the driving pressure is only dependent on P Df. During normal grain growth the pinning pressure is caused by small precipitates hindering grain boundary movement. The grain boundary is considered to be an interface between two grains of around 0.5 nm. In the case of Zener pinning, forces acting on particles and boundaries are directly transferred to each other since the whole system is in a solid state. If the grain boundary is replaced by a liquid film, these forces will no longer be transferred. Especially in the case where pinning particles are completely embedded in the liquid film. Now the retarding pressure is no longer part of equation 6.1. This explains why a migration of a liquid film is observed in strained samples that do not recrystallize before the solidus temperature of the clad alloy is reached. In the case of SILFM the driving force is: 2 α( ρf ρb) Gb 2γb P= ( PDf PDb) + PC = + or 2 R 2 αρfgb 2γb P= PDf + PC = + (6.2) 2 R This equation would suggest that when a liquid wets the grain boundaries although no strain is applied; a driving pressure would be present causing the liquid film to move. When no strain is applied the driving pressure would be: 2γb P= PC = (6.3) R This equation has been deduced by Burke and Turnbull [3] for the kinetics of grain growth. Equation 6.3 also provides the driving pressure for the process of Liquid Phase Sintering [2]. In the present study, pre-straining the materials before brazing is a pre-requisite for the onset of LFM. The driving pressure is formulated by equation

108 The grain boundary moves at a velocity (v) in response to the driving pressure (P). The velocity is assumed to be proportional to the pressure and can be represented by equation 6.4 [4] v= M P (6.4) where M is the mobility. Annavarapu and Doherty [2] deduced the velocity of the liquid film wetted grain boundary 2γ with the driving force being the capillary pressure represented by P = b. ρ v= D 4 V γ X l δ ρ R T( X l l (1 X X s ) l 2 ) where δ is the thickness of the liquid film and ρ is the grain size. (6.5) If the capillary pressure is introduced in equation 6.4 then the mobilitym is: Dl 2 V Xl(1 X M = δ R T( X X ) l s l 2 ) In this case, both the reduction in grain size and the residual strain present in the material at the brazing temperature are the driving forces for the movement of the liquid film. The overall relation for the velocity of the liquid film is found by combining equation 6.6 and 6.2 resulting in: (6.6) v= D 2 (1 ) 2 Gb l V Xl Xl γ αρf + 2 δ R T( X X ) l s ρ 2 2 (6.7) In equation 6.7 most of the parameters can be measured, with one exception namelyρ, the dislocation density in the grain just before the liquid film passes through. Quantifying the amount of dislocations present at the moment the liquid film passes through the grain is very difficult. First of all the amount of dislocations introduced by stretching the materials before brazing is already low and secondly during the brazing cycle recovery will take place, thereby reducing the number of dislocations by annihilation or rearrangement. The typical dislocation content of stress-free pure aluminium is of the order of m -2 [4]. Based on the applied strain, several authors [5,6,7] have published values for the dislocation density in aluminium alloys on a theoretical basis. The number of dislocations introduced is dependent on the type of alloy. It is widely known that the addition of elements such as Mg and Mn create strength in aluminium since they form obstacles thereby preventing dislocation movement. A major drawback of these model calculations is that the applied models are primarily focused on predicting dislocation densities at higher strains ( ε >0.1) than encountered in this thesis. Secondly, to calculate the liquid film velocity in equation 6.7 the dislocation f 100

109 density at the moment is needed when the liquid clad alloy starts to penetrate the grain boundaries, while most models just predict the room temperature dislocation density before any high temperature treatment. Most models are developed to predict the onset of recrystallization or the texture development during hot and cold rolling with subsequent annealing. To have the correct input for equation 6.7, information about the number of dislocations that survive the high brazing temperatures is required. Unfortunately this information at the time of writing is not available. However equation 6.7 can be used to make an estimate of the velocity of the liquid film. The possible parameters that can vary in equation 6.7 are: 1. δ = thickness of the liquid film (m) 2. ρ = grain size diameter in the core alloy (m) 3. ρ f = dislocation density (m -2 ) The remaining parameters are: D l V X l X s = 2.099x10-9 m 2 /s, diffusion coefficient of Si in liquid aluminium = 1x10-5 m 3 /mol, molar volume aluminium = , molar concentration Si in the liquid film = , molar concentration Si in the solid behind the liquid film R = J/molK, gas constant T = 868 K, temperature of the experiments (595 C) γ = N/m, surface tension of the liquid clad alloy α = 0.5, a constant G = 25.4x10 9 Pa, shear modulus of aluminium b = 0.286x10-9 m, Burgers vector The following graphs show the dependency of the velocity on the grain size. For the dislocation density the ranges as calculated by Hansen [5], Wang et al. [6] and Goerdeler [7] were taken. Any loss of dislocations through recovery was not taken into account. Typical values for both the grain size and film thickness were used. The thickness of the liquid film used in the first graph to calculate the dependency of the velocity on both dislocation density and grain size is 2 µm. From figure 6.2 it is evident that small grains have a strong effect on the velocity of the film, which is understandable since there is more energy per volume available in the system. Figure 6.3 is again a representation of equation 6.7 but with a grain size of 100 µm and with increasing film thickness. 101

110 Fig. 6.2: Liquid film velocity according to equation 6.15 with a film thickness of 2 µm. Fig. 6.3: Liquid film velocity according to equation 6.15 with a grain size of 100 µm. In this study for the non homogenised material with 10% stretch, the velocity of the liquid film is represented by: v= t 0.5 x10-6 (m/s) The average grain size is in the range of µm. The boxes in both figures 6.2 and 6.3 represent the velocity as calculated using the above formula in the time range between 0 and 180 seconds. From figures 6.2 and 6.3 it is clear 102

111 that equation 6.7 produces velocities that are in the range of the velocity measured in this study. The major flaw of equation 6.7 is that the liquid film velocity is independent of time, while the measurements have shown a dependency on time. Most probably the dependency on time originates from increase of the liquid film thickness and/or decrease of dislocation density with time. Both factors will decrease the velocity of the liquid film. However, due to the high temperature of the experiment (595 C) it is expected that the most of the dislocations in the material are recovered to form a stabilized structure. Recovery kinetics are strongly temperature dependent. Extrapolation of the results presented by Sandström et al. [8] for the recovery of a 1 wt% Mn alloy would result in recovery times of the order of seconds. Most likely the thickness of the film will increase with time. Annavarrapu and Doherty [2] attribute the increase in film thickness to the increase of liquid fraction during liquid phase sintering. In this thesis the amount of liquid phase stays unaltered. At 595 C, a 10 wt% Si clad alloy is completely molten. The only liquid available comes from the molten clad alloy, which remains fairly constant during the time of the experiments. Some of the molten liquid clad alloy is consumed by the affected area, since the silicon level in solid solution behind the liquid film is 1.2 wt% while in front of the film it is 0.1 wt%. Any depletion of silicon from the liquid film is replenished from the liquid clad alloy as presented in chapter 3. This silicon has to be transported from the liquid clad alloy towards the liquid film. Figure 6.4 shows the paths for silicon transportation. Fig.6.4: Silicon transport from the clad alloy to the liquid film. 103

112 The relevant question whether silicon is transported by diffusion or convection from the clad alloy towards the liquid film has been answered by means of the following experiment. Aluminium brazing can be done by using a high vacuum or a controlled atmosphere with flux. In a niche market a third process is being used called fluxless brazing. Fluxless brazing is facilitated by first electrochemically plating the clad alloy with a nickel-bismuth alloy [9,10]. This nickel-bismuth alloy reacts exothermic with the clad alloy, thereby disrupting the oxide layer making it possible for the liquid clad alloy to flow and form joints. The actual mechanism has been described in more detail by Cheadlle and Dockus [11]. Nickel is almost insoluble in solid aluminium (0.05 wt%) while it has a good solubility in liquid aluminium (6.12 wt%). During the time where a liquid clad alloy is available all nickel is in the liquid phase as a ternary eutectic [11]. If transportation from silicon is taking place through the liquid phase, nickel should also be present in this liquid. If a brazing sheet alloy susceptible to LFM is brazed with a nickel plated clad alloy, the position where after brazing nickel is found should give a clue to how silicon is transported from the clad to the liquid film. For this experiment, a production modified 3003 core alloy, double side clad was nickel plated and brazed at 595 C. Cross sections were made and analysed by OM and SEM with EDX for the element distribution. Figure 6.5 shows the OM image of the 4% stretched sample after brazing. Again the liquid film and affected area can be seen. Fig. 6.5: OM image after 4% stretching and subsequent brazing of Ni plated brazing sheet. Figure 6.6 shows the SEM image and a Back Scatter Electron image from an area with a connection between the clad alloy and liquid film. 104

113 Fig. 6.6: SEM and BSE image from the affected area. From the detailed section of figure 6.6 an element mapping was made. Figure 6.7 shows the results. It is evident that the nickel is located in the liquid film but also in the channel feeding the liquid film from the residual clad. If the silicon in the liquid film was fed by solid state diffusion from the clad alloy, no nickel would show up due to its low solid solubility. The round sphere in the detailed image of figure 6.6 is another clue that a liquid is passing through a solid; this will be addressed in the discussion hereafter. Fig.6.7: Element distribution. Ludwig and Bellet [12], Ludwig et al. [13], Ina and Koizumi [14] and Pereiro-Lopez et al. [15], studied the kinetics of gallium penetration in aluminium. It was observed that the penetration process could be divided into three sections: 1. Incubation: It takes some time to detect the first traces of gallium in the grain boundaries. 2. Penetration: Formation and thickening of a liquid film between the grain boundaries. 3. Saturation: Thickening of the film stops. It is assumed that the penetration of the molten silicon clad alloy shows a similar behaviour, the observations made in chapter 4 can be explained accordingly. Figure 4.5 shows that it takes some time (incubation) to get the liquid film to move and to reach its maximum velocity. 105

114 After reaching a critical thickness, the velocity decrease again due to the thickening of the liquid film (penetration), which is in line with equation 6.7. When equation 6.7 is used and the grain size is taken to 50 µm and the dislocation density to be 1x /m 2 the relationship between liquid film thickness and velocity is as illustrated in figure E-06 liquid film velocity (m/s) 4.00E E E E E liquid film thickness (µm) Fig. 6.8: Relation between film thickness and film velocity. Figure 6.8 shows a hyperbolic relation between thickness and velocity, which correlates better with the actual observation. How the film thickness evolves in time is still uncertain. Wolski and Laporte [16] reported a liquid film thickness dependence of 0.5 t for the nanometric film becoming a at micrometric thickness for nickel grains wetted by liquid bismuth. Assume the thickness evolution to be dependent on: 0.5 δ =Ca t (6.8) 0.33 t dependence Ca is a constant but it remains unclear which physical parameters determine the size of the constant. If equation 6.7 is used and all parameters are considered to be constant except for the film thickness, than combination of equation 6.7 and 6.8 become: 0.5 v = C t (6.9) D 2 (1 ) In which 2 Gb l V Xl Xl γ αρf C = + 2 Ca R T( Xl Xs) ρ 2 Their observations are in line with the observations of this study. Although there seems to be some agreement on the film thickness evolution in time, it is still not clear which physical parameters play a role in the film thickness development

115 Possibly, the increase in the liquid film thickness is just based on fluid mechanics. Behind the moving liquid film silicon is being taken up by the matrix to reach its equilibrium concentration of 1.2 wt%. This will drain the silicon from the liquid film and is then replenished from the remaining clad alloy as can be seen in figure 6.4. The precipitates swept from the affected area will also be concentrated in the liquid film. To allow the film to move the precipitates have to be separated by a distance allowing the liquid to pass through. This introduces another variable into the system, making a prediction of the film thickness development in time even more difficult. 6.3 Discussion Alternative models The movement of the liquid film is closely related to the melting or dissolution at the front and solidification in the wake of the liquid film. In principal the liquid film consists of two moving solid-liquid interfaces. These interfaces move at almost equal velocity. In case the film thickness increases in time there is a small difference in speed between the front and the back of the liquid film. In the past the movement of the liquid film has been formulated either by applying theoretical considerations or just by fitting the experimental observations. In this study the best fit for the velocity and travelled distance were presented in chapter 4 and are repeated here to make an easier comparison with the results from other authors. v= t S 0.5 x10-6 (m/s) 0.5 = t x10-6 (m) Some of the theoretical formulations for the velocity of or distance travelled by the liquid film are: C Dl v= C(1 k) δ l [17] (6.8) Where Cis the concentration difference across the liquid film, C l is the concentration in the liquid film at the trailing interface, k is the solute distribution coefficient, D l is the diffusion coefficient for the solute in the liquid and δ is the thickness of the liquid film. According to Brechet and Purdy [17] the driving force for liquid film migration was considered to be the coherency strain as discussed in chapter The proposed model of Brechet and Purdy [17] is solely based on the assumption that the coherency strain generated by diffusion in front of the liquid film is the driving force, while in our study the reduction of grain size diameter and the residual strain present are considered to be responsible for the movement of the liquid film. 107

116 Fredriksson et al. [18] and Antonsson et al. [19] attributed the movement of the liquid film to grain boundary wetting as discussed in chapter ~ l ( CA Ce) Dl v= [18,19] (6.9) l s λ ( C C ) e e The first term represents the concentration difference across the liquid film, D l is the l s diffusion coefficient in the liquid and ( C ) is the difference in solute concentration e Ce between the unstressed solidified solid and liquid at the trailing side of the liquid film. Yoon [20] attributes that the movement of the liquid film is caused by the release of the coherency strain as presented in chapter CV 0 D S = mγ [20] (6.10) RTdt S is the distance travelled by the liquid film. C0is the equilibrium solubility of Si, V m is the molar volume of the solid, γ is the surface tension of the liquid, D is the diffusion coefficient in the liquid, d is the grain size, t is the thickness of the liquid film, R is the gas constant and T is the temperature. In this study coherency strain does not relevant, therefore equation 6.10 is not considered to be valid for this study. Nylen et al. [21] suggested that the reduction in grain size is responsible for liquid film migration as discussed in chapter and again there is no reference to any strain component. Hillert [22] and Kuo and Fournelle [23] both claim that the driving force for LFM is not the coherency strain energy but the free chemical energy change. However, if this were valid, liquid film migration would occur in all recrystallized samples from this study, which is not the case. Brener and Temkin [24, 25, 26] have developed a theory to describe the movement of two solid-liquid interfaces. One is the melting front and the other is the solidification front. The interfaces interact through diffusion in the liquid layer between the two interfaces. In their description, the coherency strain caused by the diffusion in front of the liquid film is responsible for the movement of this liquid film. Their treatise solving the mathematical problem is complex but the resultant equation to describe the liquid film velocity is relatively simple. 2 3 Db v [24,25,26] (6.11) d 0 D is the diffusion coefficient of the solute in the liquid film, b is a measure of the size of the coherency strain, is the term which describes the concentration difference across the liquid film between the melting and solidification front and d0 corresponds with the thickness of the 108

117 liquid film that is in equilibrium with the melt. Again this model does not take into consideration the residual strain left in the material at the brazing temperature. Schatt at al. [27,28] mentioned the influence of the dislocation density on solid state and liquid phase sintering. The migration velocity was only dependent on the dislocation density difference across the liquid film. No reference was made to the importance of the size of the particles that were sintered in their study. It was mentioned that the cold milled powder with a high dislocation density showed more LFM than the non-cold milled powder. As presented in this section, there are a number of models describing the movement of the liquid film. Basically, the models presented can be divided into two groups, one group that assumes that the coherency strain is the driving force and the other group that assumes that the reduction of the surface energy is the driving force for LFM. However, all these equations lack a term that incorporates also the contribution of the strain as a possible driving force for liquid film migration. Equation 6.7 does not show a dependency on time, whereas the relation between velocity and time was found to be v= t 0.5 x10-6 (m/s). The remaining dislocation density at the brazing temperature is a function of time and temperature. High temperatures and long holding times at high temperature will decrease the dislocation density [4]. The extent to which recovery will take place is dependent on the alloy composition and thermo-mechanical processing prior to brazing. Fukumoto and Doko [29] clearly demonstrated that the type of annealing had an effect on the recrystallization of the core alloy resulting in the occurrence of LFM. A batch type anneal with a slow heat up rate resulted in a much larger grain size compared to the rapidly heated continuously annealed material. Based on equation 6.7, materials with small grains would show more liquid film migration compared to coarse-grained materials at similar dislocation densities. The photos in figures 6.9 are taken from [29] and show the effect of grain size on the occurrence of LFM for a clad AA3003 alloy. Continuously annealed 0% 3% 5% 109

118 Batch annealed 0% 3% 5% Fig. 6.9: LFM in two different grain sized materials at different stretch levels. The continuously annealed material has a much smaller grain size than the batch annealed material, therefore even at zero applied strain the fine grained material will show LFM. The coarse-grained material starts to show LFM when some strain is applied and the core alloy does not recrystallize. The fine-grained material starts to recrystallize at lower strains than the coarse-grained material. The effect of grain size on LFM was also reported by Wittebrood et al. [30]. In figure 6.9 the LFM depth in both materials is the same only the maximum penetration depth is dependent on the level of stretching applied prior to brazing. The effect of strain on the occurrence of LFM has also been found in two other metallurgical systems. Schatt et al. [27, 28] showed that the introduction of dislocations by ball milling had a significant effect on the liquid phase sintering of tungsten powder. The ball milled tungsten powder showed much more penetration by the liquid nickel during sintering causing a chemical change in the tungsten powder. A similar observation was made by Antonsson et al. [19], where cold worked tungsten rod showed much deeper penetration by a nickel iron melt than a fully recrystallized tungsten rod. Another feature that can be found in the literature related to a moving liquid film is the detachment of pockets filled with precipitates or constituents from this liquid film [2,31,32]. Somehow, the drag of these clusters of particles is so large that they cannot be carried any more by the moving liquid film. Figure 6.10 shows an example of the partial and complete detachment of a cluster of particles observed in samples from this thesis. 110

119 Fig.6.10: Partial and complete detachment of particles. Figure 6.9 shows great resemblance to the series of schematic drawings made by McPee et al. [32], as presented in figure The left hand micrograph from figure 6.10 would correspond to his suggested stage 4 while the right micrograph resembles stage 5 in figure Figure 6.10 also shows a detachment of Mn-bearing precipitates embedded in a solidified pool of ternary eutectic liquid. Fig. 6.11: Different stages of the detachment of particles. [32] Another observation supporting the importance of dislocations or strain energy as a driving force for LFM can be seen in figure 6.12 where an area with localised LFM is shown, Fig. 6.12: Localised LFM. The material in figure 6.12 has been stretched 3% just prior to brazing. 111

120 It is well known from the literature that deformation rarely occurs in a homogeneous manner. In particular, single crystals with special orientations and coarse-grained materials show very heterogeneous deformation behaviour [4]. At low strains (ε <0.1) not all grains will deform in the same way. The deformation behaviour of the grains will depend on the crystallographic orientation of the individual grains with respect to the direction of deformation. Grains with a specific orientation will deform more at low strains, meaning their dislocation density will be higher compared to adjacent grains. The uneven distribution of dislocations can possibly result in local areas in the material where SILFM can occur; while in the same material the majority of the structure does not show any SILFM as is illustrated in figure Chen [33] demonstrated that the recrystallization kinetics for different texture components in aluminium after the same applied strain can be completely different. Brass, copper and S texture components recrystallize easier than the cube component. Abnormal Grain Growth occurs in materials with unstable microstructures resulting in the excessive growth of a few grains. AGG is sometimes known as secondary recrystallization [4]. The driving force for AGG is the local reduction of grain boundary energy. Several publications on modelling AGG suggest that AGG can occur if there is an uneven distribution of grain boundary surface energies caused by wetting [34, 35, 36]. In this case only one side of the grain is wetted, being the side in contact with the molten clad alloy. In this case the criteria for an uneven distribution of grain boundary energy as presented by Hwang et al. [35] is fulfilled and grain growth could occur. The affected areas through which the liquid film has changed the composition can be seen as large grains. Figure 6.13 shows that SILFM results in huge grains behind the liquid film. AGG shows similarities with SILFM. Fig. 6.13: Abnormal grain growth as a result of SILFM. 112