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本科毕业设计(论文)外文翻译（附外文原文） 学 院： 机械与控制工程学院 课题名称： 小皮带线对象系统远程 监测技术研究 专业(方向)： 机械设计制造及其自动化 （模具设计与制造） 班 级： 机械11-2班 学 生： 谭振彬 指导教师： 周轶旻 日 期： 2015年3月25日 外文翻译先进的高强度沉淀硬化钢用于汽车的设计和应用开发摘要严格的安全和污染控制规范迫使汽车零部件使用高强度钢。然而，高强度钢通常具有较低的可成形性，这使得成形操作困难。因此，开发高强度钢以及汽车应用更好的成形性，是所有钢铁企业面对的一个挑战。在目前的工作中，一种新的方法，从钢的混合物、设计参数到开发高强度钢（最小UTS 1000MPa）延伸率高（最低15%）进行了详细的论述。这种钢的高延伸率有望让成形操作变得简单。关键词：合金设计，热轧，连续降温，高强度，高延伸率，晶粒大小，析出强化。1 简介车体的不同组分需要许多复杂的形成操作，其中，所述可延伸材料是必不可少的。汽车制造商都在增加使用高强度钢，以符合不同区域的安全法规，同时也降低了汽车的重量，使燃料更充分的利用。这从汽车行业持续增长已迫使钢铁行业推到达了一个新的高度。不同的钢铁制造商已经开发出不同的路线和理念（如双相或双相钢，相变诱发塑性或TRIP钢等），汽车行业可根据自己的条件来满足这种需求。虽然采用先进的高强度钢（AHSS）在过去的几年中显著增加1,2，由于这些钢具有模塑性差，汽车零部件复杂的成型操作使得这些钢广泛应用于汽车行业有所困难3。这有必要建立高强度钢（最小UTS 1000MPa）高延伸率（至少15），以促进良好的成形操作。最近，塔塔钢铁公司研究人员4,5已经证明成功的厂沉淀硬化的高强度钢至少要600和800MPa UTS。在这些发展情况下，增加轧制不在热轧、冷却和卷取也是没有问题的。这两种钢现在正在工厂生产，轧制力，冷却速度和卷取机仅适用于一般无间隙原子（IF）钢。因此，我们决定延长了同样的想法，开发了一个新的档次高强度钢（UTS 1000MPa的），它可以热轧，冷却和卷取容易。考虑冶金，合金设计，显微结构和工艺参数的选择最终在钢声学特性的机械性能和新开发的在目前的工作介绍。2 理论设计和热力学分析合金成分设计要记住这些限制。因为它会促进透明的铁蛋白微合金转化为微细碳化沉淀物。热轧带卷从奥氏体区开始降温到完全冷却的这段时间里，带卷的上表可能会出现这种沉淀。因此这种钢的屈服强度是由于合金元素的固溶强化,显微组织、晶粒细化(粒径)、位错和沉淀硬化。自从生产钢板热轧以来,从来没有考虑过位错硬化。钢的屈服强度可以通过下列方程表示6:y=o+ ss+gs+ppt （6）其中，o表示铁的强度（兆帕），ss表示固溶强度（兆帕），gs表示晶粒强度（MPa）和ppt表示沉淀强度（兆帕）。碳是碳化物形成钢的一种非常有效的强化剂，它取决于微合金化元素，如钛、钼、铌、钒7。为了达到抗拉强度1000MPa,碳含量应该在0.060.15。低碳钢在高温形成钢是会产生渗碳体和珠光体，这些成分对钢的韧性是非常有害的。因此应该限制他们的形成。因此,应该添加锰降低奥氏体向铁素体转变温度(Ar3)8来限制这类粒子的形成。虽然硅对固体融化很有效 9,但是为了减少硅的产生,其用量只含0.2 。设计的目标是开发一个沉淀强化钢,因此微合金化在合金设计中扮演着重要角色。钛和钼中加入作为析出强化元素，提高了每单位体积10析出物的数量。有人建议早期Ti含量应为0.1-0.35和Mo小于0.25（以重量计），为了达到约100MPa的强度11还添加一定量的铌和钒，利用其拉伸强度的优点形成复杂的沉淀组织钛和钼 12。因此，可以预期的是任何短程1000MPa的放置目标将由加法得到补偿。然而，当的浓度变得高于0.08或增加的钒的浓度超过0.15，钢的伸长会降低11。另外，钛和钼的原子保持在比任何降解（钛/钼）90的铁素体）是只用在第二阶段，并且可以在微观结构中可以看出铁素体以体积分数的形式是有一些差异。保持的温度越高,对第二相粒子的伸长率越不好。因此，为了限制第二相粒子的生成，卷取温度应保持低温。但是，随着卷取温度的下降铁素体也会相应的增加。从图1中可以看出，理论上钢的温度在570590时开始生成贝氏体的。另一方面，处于更高的卷取温度时，主要的结构是多边形铁素体。因此，应处于的CT的最佳温度范围，从而限制第二相粒子的生成，并且生成最多的多边形铁素体。硬度测量的模拟样本如图7所示，这清楚地表明，模拟卷取在低温范围内来获得最大硬度（在600615下的维氏硬度为250）。这个硬度值预计可以提供850900MPa的拉伸强度16。虽然这比目标拉伸强度略低，但它决定热轧后探索的可能性。因此，卷取温度决定维持在61515，并相应地进行了进一步的热轧实验。4.3 实验热轧实验室仿真热轧和卷取已经在图3和前一节中描述。图8是在光学显微镜下热轧的微观结构。从图8中可以看出, 在光学显微镜中观察每个试样在2.73.4毫米的范围内的10个不同的位置，没有看见出第二相的铁蛋白结构（如珠光体/贝氏体/马氏体）。为了更好地观察微观结构,利用扫描电子显微镜(SEM)观察到的微观结如图9所示。图9的a和b显示第二阶段没有以铁素体为主的显微组织。只考虑大倾角晶界（4151），扫描超过1500粒电子背散射衍射（EBSD）的分布粒子为0.2mm。EBSD分析如图10和图11所示，图10表示扫描区的对比度(BC)，图11表示整个区域粒度分布百分比。可以观察到几个较大的晶粒和大量的小颗粒微观结构。从图11中可以看出，尽管晶粒的分布是不均匀的，但大部分的颗粒小于3毫米，它被认为是非常完善尺寸分布图。这种细颗粒的分布可以改善钢的力学性能。这种钢可以从纳米析出物中得到很好的性能。为了识别和描述这些析出物，使用透射电子显微镜(TEM)进行了详细的研究，图12和图13是合金1和合金2分别透射的电子显微照片。可以看到大量的超细硬质合金析出物的直径为几个纳米大小(大约24纳米)。用X射线光谱仪（EDX）分析这些析出物的色散，识别出复杂的碳化物析出物（钛、钼和铌）。比较图12和图13可以看出，合金2中析出物的数量远远多于合金1，这可能是合金2存在较高含量的碳和金属元素如钒。钢的力学性能通过单轴拉伸试验的结果如图14所示。从拉伸试验结果可以看出，虽然合金1的力学性能不如合金2,但两个合金所需的最小强度都是1000Mpa。然而, 正如前面提到的，钢的力学性能受微观结构的影响，该合金的强度可以在许多强化机制中发挥重要作用（等式（1）。这里不考虑钢这种简单材料热轧后位错的强化和提高。考虑固溶体硬化，在皮克林不同的基础上都以公式（2）来表示6。在计算过程中，碳影响铁素体的最大溶解度，由于部分碳结合钛和钼，因此不利于固溶强化。这样计算一个固溶的硬度大约为120Mpa。ss=5544wt%C+5544wt%N+678wt%P+83WT%Si+32wt%Mn+11wt%Mo （2）从图11a和b可以断定，大约60颗粒的直径大于3毫米。因此，平均粒径3毫米是由霍尔方程计算出来的。公式（3）是计算小颗粒的尺寸。gs= KD-0.5 （3）其中，k为常数，并且平均晶粒直径的常数值如前面所述17。用这个公式计算，可以知道细小晶粒的分布会导致合金的屈服强度为420450MPa。因此，考虑到铁的强度为56MPa，从这两种合金中析出的粒子的直径在24纳米范围内，他的屈服强度为330430MPa13,18。从图12和图13观察合金析出物的屈服强度，可能是由于下列原因：首先，合金2的析出物比合金1多，其次，合金2含有一下钒，这可能会导致合金2的析出物有更强的屈服强度。5 总结全铁蛋白钢的设计理念有助于增加硬质合金的析出物。确定钢的成分要考虑不同来源的和记住植物功能，测试机模拟测定热轧的工艺参数(完成轧制和卷取温度)，最终以微观结构来推理力学性能。上面的工作总结如下几点：（）高强度钢结合了很好伸长率。（）硬质合金钢可以增加不同合金元素的固溶体析出铁素体的数量。（）力学性能实现的最低要求如下：屈服强度(YS)为89Mpa；极限抗拉强度(生产资料)为1020Mpa；14%的均匀伸长和总延伸率为18%。（）固溶强化的作用大约为120Mpa。（）细的铁素体晶粒尺寸(3毫米)。这样的铁素体组织可以增加的强度为420450Mpa。（）纳米合金碳化物的析出物可以增强330430Mpa的力学性能。致谢作者想感谢塔塔钢铁有限公司的管理层批准著作的出版。还要感谢国家冶金实验室进行詹谢普尔一些重要的测试。外文原文Design and development of precipitate strengthened advanced high strength steel for automotive applicationAbstract: Stringent safety and pollution control norms forces the car industry to use higher strength steels indifferent components of an automobile. However, high strength steels typically possess less formability which makes forming operations difcult. Thus developing high strength steel, with better formability for automotive application, is a challenge for all the steel companies those are working to solve indifferent ways. In the present work, a new approach, from designing of steel composition to nalizing the process parameters, to develop high strength steel (minimum 1000MPa UTS) with high elongation (minimum 15%) will be discussed in details. High elongation of this steel is expected to make the forming operation easier without compromising with the strength.Article history: Received 18 August 2012, Received in revised form11 October 2012, Accepted 13 October 2012, Available online 23 October 2012Keywords: Alloy design, Hot rolling, Continuous cooling, High strength high elongation, Grain size, Precipitation strengthened.1. IntroductionDifferent components of a car body require many complicated forming operations where ductility of the material is very crucial. The automotive manufacturers are also continuously increasing the use of high strength steel in order to comply to the safety norms of different geographic regions and also to reduce the weight of the car to make it more fuel efcient. This continuous thrive from the automotive industry had forced the steel industry to push its horizon to a never achieved region. Different steel producers have developed different routes and philosophy (e.g. Dual Phase or DP steel, Transformation Induced Plasticity or TRIP steel etc.) depending on their mill condition to meet this demand from the automotive industry. Although the use of advanced high strength steels (AHSS) has been increased signicantly in the past few years 1,2, as these steels have poor formability, it is very difcult to apply these steels in a wide range of automotive components due to complicated forming operations 3.This has necessitated the need to develop high strength steel (minimum 1000MPa UTS) with high elongation (minimum 15%) to promote good forming operation to be carried out. Recently, researchers 4,5 at Tata Steel had demonstrated successful plant production of precipitation hardened high strength steel grades with UTS of minimum 600 and 800MPa. In those developmental cases, roll load did not increase during hot rolling and cooling and coiling also was not a problem. Both these two steels are now being produced in a plant where the roll force, cooling rate and down coiler are suitable for producing typically Interstitial Free (IF) steels only. Thus, it was decided to extend the same philosophy to develop a newer grade of high strength steel (UTS1000MPa) which can be hot rolled, cooled and coiled very easily. The metallurgical considerations, design of alloy, selection of processing parameters and nal characterization of micro-structural and mechanical properties of this newly developed steel will be described in the current work.2. Theoretical design and thermo dynamical analysisThe alloy composition was designed keeping in mind the restrictions mentioned above. Thus, it was decided to have a completely ferritin microstructure dispersed with interphase precipitates of micro alloyed carbides. This precipitation can take place during cooling on run out table or after coiling when the coil starts to cool naturally after hot rolling from the austenite region. Thus the yield strength of this steel becomes a function of solid solution strengthening due to alloying elements, microstructure, grain renement (grain size), dislocation and precipitation hardening. Since, the steel sheet is produced in hot-rolled condition, the contribution from dislocation hardening was not considered. The yield strength of steel can be expressed by the following equation 6: y=o+ss+gs+ppt (6) Where, o is strength of pure iron (MPa), ss is the solid solution strengthening (MPa), gs is the strengthening due to grain size renement (MPa) and ppt is the strengthening due to precipitation (MPa). Carbon forms carbides and is one of the very effective strengtheners of steel depending on the contents of microalloying elements e.g. Ti, Mo, Nb and V. It was shown earlier 7 that in order to achieve tensile strength around 1000MPa, the carbon content should be in the range of 0.060.15wt%. Coarse cementite and pearlite can be formed in low carbon steel if such steels are coiled at high temperature. These constituents are very much detrimental for the toughness of steel. Thus their formation should be restricted. Therefore, manganese should be added to lower the austenite to ferrite transformation temperature (Ar3) 8 in order to restrict the growth of such particles. Although Si is very effective for solid solution strengthening 9, to minimize the occurrence of Si-scale, its usage was limited to only 0.2wt%. As the objective here was to develop a precipitate strengthened steel, selection of precipitates and thus the microalloying addition plays a key role in alloy design. Titanium and molybdenum were added as precipitation strengthening elements to increase the number of precipitates per unit volume 10. It was suggested earlier that the Ti content should be within 0.10.35wt% and Mo should be more than 0.25wt% in order to attain tensile strength around 100MPa 11.Some amount of Nb and V were also added to utilize their ability to rene the microstructure and to form composite precipitates together with Ti and Mo 12. Thus, it was expected that any short fall in reaching 1000MPa target would be compensated by such addition. However, when the concentration of Nb becomes higher than 0.08wt% or the concentration of V increases above 0.15%, elongation of the steel gets deteriorated 11. Moreover, the atomic concentration ratio of Ti to Mo was kept below (Ti/Mo) 90% ferritic) only with some difference in volume fraction of the second phase and morphology of ferrite. With the higher holding temperatures, the second phase particles become coarser which is not good for elongation. Thus, to restrict the growth of such second phase particles, the coiling temperature should be kept to the lower side. However, acicularity of ferrite increases with lowering the coiling temperature. This is expected as the theoretical bainite start temperature of this steel is in the range of 570590, as can be seen from Fig.1. On the other hand, for higher coiling temperatures, the structure consists mostly with polygonal ferrite. Thus, an optimum range of CT should be used so that the growth of the second phase particles can be restricted as well as the ferrite remain in polygonal form to the maximum possible extent.Hardness measurements of such simulated samples are shown in Fig.7, which clearly reveals that the maximum hardness is obtained at the lower range of simulated coiling temperature (250 VHN at 600615). This hardness value is expected to deliver a tensile strength of 850900MPa 16. Although this expectation is somewhat less than the target property, yet it was decided to explore the possibilities after hot rolling. Accordingly the coiling temperature was decided to keep at 61515 and further experimental hot rolling was carried out accordingly.4.3. Experimental hot rollingSchematic of the laboratory scale hot rolling and coiling simulation is shown in Fig.3 and the procedure has been described in the previous section. Fig.8 represents the hot rolled microstructure under optical microscope. It can be observed from Fig.8 that a predominantly ferritin microstructure was produced and presence of hard second phase (e.g. pearlite/bainite/martensite) was not detected. The average grain size from optical microstructures calculated using 10 different locations of each samples was in the range of 2.73.4 mm.In order to reveal the microstructures in a detailed manner, scanning electron microscope (SEM) was utilized and the microstructures are shown in Fig.9. Both Fig.9a and b show a predominantly ferritic microstructure with apparently no presence of any second phase. The grain size distribution was carried out through electron back-scattered diffraction (EBSD) with a step size of 0.2 mm considering only the high angle grain boundaries (4151). More than 1500 grains were scanned and Figs.10 and 11 represent the EBSD analysis where Fig.10 shows the Band Contrast (BC) image of the area scanned and Fig.11 reveals the percentage wise grain size distribution for the entire area. It can be observed that a few larger grains along with much higher number of smaller grains are present in the microstructure. Although heterogeneous distribution of grains is apparent, yet it can be seen from Fig.11aandb that most of the grains are less than 3 mm in size which can be considered as very ne scale distribution. This very ne distribution of grains can lead to signicant strength improvement in the nal steel.This steel was designed in such a way that it can also derive signicant fraction of its strength from the ultra-ne, nano-sized precipitates. In order to identify and characterize those precipitates, detailed study was performed using transmission electron microscope (TEM). Figs.12 and 13 show a series of transmission electron micrograph image of both Alloys-1 and -2, respectively. A large number of ultrane carbide precipitates with a size of few nanometers in diameter (approximately 24 nm) can be seen. Energy Dispersive X-ray Spectroscopy (EDX) analysis performed on these precipitates particles identied them as complex precipitates of (Ti, Mo, Nb) Carbides. A comparison between Figs.12 and 13 reveals that the number of precipitates in alloy 2 is much higher than alloy 1. This may be attributed to the higher content of carbon as well as the presence of additional microalloying element such as vanadium.Final mechanical properties were evaluated by performing uniaxial tensile test and the results are shown in Fig.14 for both the steels. From the tensile test results it can be seen that both the alloys are meeting the minimum required strength of 1000MPa though, Alloy 1 is inferior to the Alloy 2 in terms of strength. However, it is worth to discuss the origin of such outstanding strength in these steels in details. Many strengthening mechanism, as mentioned earlier, can play important role in developing the strength of an alloy (Eq. (1). As the steel was hot rolled and single phase material, dislocation strengthening and strengthening due to second phase particles were not considered here. For considering the solid solution hardening, contributions of different elements were considered as described in Eqn.(2) after Pickering 6. During the calculation, maximum solubility of carbon in ferrite was considered because rest of the carbon would be bonded with Ti and Mo as precipitate particles and thus would not contribute to solid solution hardening. Such a calculation yielded a possible contribution from solid solution hardening to about 120MPa.ss=5544wt%C+5544wt%N+678wt%P+83WT%Si+32wt%Mn+11wt%Mo (2)From Fig.11a and b, it can be concluded that approximately 60% of the total grains examined through EBSD are less than 3 mm in diameter. Thus, the average grain size was considered as 3mm and Hall-Petch relationship was applied as described in Eq. (3) to calculate the contribution originating from ne scale grain distribution.gs=KD-0.5 (3)Where k is a constant and the average grain diameter is D. The value of k was considered as mentioned by Takeda et al. 17. From this calculation, it can be expected that ne grain size distribution can result approximately 420450MPa contribution towards the nal yield strength of the alloy. Thus, considering the base strength of iron as 56MPa, it was apparent that the contribution from the precipitate particles in these two alloys can be in the range of 330430MPa which can be expected from earlier works also for a precipitate size 24 nm 13, 18. Such a precipitate size can also be observed from Figs.12 and 13. The higher yield strength of Alloy 2 could due to be the following reasons: rstly, Alloy 2 may contain ner precipitates than the Alloy 1 and secondly, the Alloy 2 contains some V which might hav
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