钢热处理的冶金原理

钢材热处理中的冶金原理

钢的热处理是为了实现钢的金相组织性能的预期变化。通过热处理，钢的性能会发生剧烈变化。当钢被加热到高温奥氏体状态，然后在接近平衡的条件下缓慢冷却时，通常会获得非常稳定的钢结构。这种类型的热处理，通常称为退火或正火，产生的结构具有低水平的残余应力锁定在钢中，并且可以从 Fe（铁）-C（碳）平衡图预测结构。然而，钢中最需要的特性是高强度和硬度，这些特性通常伴随着高水平的残余应力。这是由于奥氏体状态的非平衡冷却或淬火产生的亚稳态结构。

晶体结构和相

已知固态纯铁的晶体结构以两种同素异形体状态存在。从环境温度到 910 摄氏度，Fe 具有体心立方 (bcc) 晶格，称为 α-Fe。在 910 摄氏度时，α-Fe 晶体变成具有面心立方 (fcc) 晶格的 γ-Fe 晶体。伽马晶体在高达 1400 摄氏度的温度下仍保持稳定性。高于此温度，它们再次获得称为 delta 晶体的 bcc 晶格。 δ晶体与α晶体的区别仅在于它们存在的温度区域。 Fe 有两个晶格常数，即 (i) bcc 晶格 (α-Fe, delta-Fe) 为 0.286 nm，和 (ii) fcc 晶格 (gamma-Fe) 为 0.364 nm。在低温下，α-Fe 表现出很强的铁磁特性。当它被加热到大约 770 摄氏度时，这种情况就会消失，因为晶格失去了它的铁磁自旋顺序。 Fe在770℃以上的状态称为β-Fe。顺磁性β晶体的晶格与α晶体的晶格相同。

在从一种形式到另一种形式的过程中，Fe 能够过冷。这导致加热和冷却时转变点的位置不同。差异取决于冷却速率，称为滞后。字母“c”和“r”表示转变是由于加热还是冷却。此外，α-Fe 转变为 γ-Fe 时密度的变化导致材料体积的突然变化。有时它会产生超过弹性极限的应力并导致失效。 γ-Fe的密度比α-Fe高4%左右。

铁碳平衡图

钢的结构是 Fe-C 合金，可以包含纯 C（石墨）或称为渗碳体 (Fe3C) 的化合物作为富 C 成分。即使在相对缓慢冷却的钢中也存在渗碳体（通常需要在较高温度下长时间保持才能将 Fe3C 分解为 Fe 和 C）。因此，Fe-C 平衡图经常被视为 Fe-Fe3C 平衡图。 Fe-C 图是稳定的，而 Fe-Fe3C 图是亚稳态的。结合Fe-C稳定图和Fe-Fe3C亚稳图的Fe-C平衡图如图1所示。虚线表示稳定Fe-C图，实线表示亚稳Fe-Fe3C图。

图1铁碳图

在亚稳态 Fe-Fe3C 图中，Fe 的同素异形体晶格（δ、γ 和 α）是 δ、γ 和 C 在 Fe 中固溶体的形成位点。当贫 C 钢结晶时，δ 固溶体晶体在液相线 AB 和固相线 AH 处沉淀。 δ固溶体具有体心立方晶格。在 1490 摄氏度的最高温度下，δ 溶液含有 0.1 % C（H 点）。在 1490 摄氏度时，饱和 delta 溶液和含 0.5% C 的液体之间发生包晶反应（B 点）。结果，形成了C在γFe中的γ固溶体。含0.18% C（I点）。

如果 C 含量高于 0.5%，γ 固溶体直接从液体中结晶（在液相线 BC 和固相线 IE）。在 1130 摄氏度时，C 在 γ Fe 中的极限溶解度接近 2.0 %（E 点）。将温度从 1130 摄氏度降低会导致 ES 线处 C 在 γ-Fe 中的溶解度降低。在 723 摄氏度时，C 的溶解度为 0.8 %（S 点）。 ES线对应于γ溶液中Fe3C的沉淀。

随着C含量的增加，γ晶格转变为α晶格的温度降低，转变发生在对应于曲线GS和GP的温度区间内。 α相沉淀曲线GS与Fe3C沉淀曲线ES相交。 S点是坐标为723℃和0.80%C的共析点。此时饱和α溶液和Fe3C沉淀同时形成共析浓度γ溶液。 α固溶体的晶格与δ固溶体的晶格相同。在723℃的共析温度下，α固溶体的含量为0.02 % C（P点）。

进一步冷却导致 C 在 α-Fe 中的溶解度降低，并且在室温下它等于一个百分比的一小部分（D 点）。当 C 含量为 2 % – 4.3 % 时，结晶开始于 BC 线处的伽马溶液沉淀。 C 含量增加到 4.3% 以上会导致 Fe3C 在线 CD 处沉淀。在 C 点超过 2.0% 的所有铁合金中，剩余初生相析出后，γ 溶液和 Fe3C 在 C 点共晶结晶，其坐标为 1130 摄氏度和 4.3% C。线 Ao 与磁从铁磁态到顺磁态的转变。

在稳定的 Fe-C 平衡图的情况下，由于非常低的冷却速率，C（石墨）可以直接从液体中结晶出来。在这种情况下，形成奥氏体和石墨的共晶混合物，而不是奥氏体和渗碳体的共晶。图 1 中的虚线表示铁石墨系统。这些线的温度高于 Fe-Fe3C 系统的线。这证实了铁石墨系统的更大稳定性和接近完全平衡。这也得到了以下事实的支持：加热含有大量 Fe3C 的高 C 钢会导致其分解，如方程式 Fe3C =3Fe + C 所示。

在中等冷却速率下，一部分钢可以根据石墨系统结晶，而另一部分根据渗碳体系统结晶。两个系统图中的相平衡线可以根据特定的冷却速率进行位移。可以看到 C 固溶体在 γ-Fe（奥氏体）中的析出线有明显的位移。因此，该图仅适用于冷却速度相对较慢的钢。

碳的影响

C 在 α-Fe 中的最大溶解度出现在 721 摄氏度，等于 0.018 % C。经过淬火，C 可以保留在 α 固溶体中，但很快就会通过老化机制开始相沉淀。在固溶体中，C 可以形成（i）均匀溶液，静态均匀的间隙分布，这是一种罕见的情况，或（ii）不均匀的溶液；在晶格结构受到干扰的地方（晶界、位错）形成团簇。后者是最可能的固溶状态。这样形成的团簇阻碍了塑性变形过程中位错的运动，并导致塑性流动开始时变形的不均匀发展。

为了分析 C 含量对 Fe-C 合金的影响，需要考虑每个结构成分。缓冷钢由铁素体和渗碳体或铁素体和石墨组成。

铁氧体是塑料的。在退火状态下，铁素体具有较大的伸长率（约 40%），柔软（布氏硬度为 65 -130，取决于晶体尺寸），并且在高达 770 摄氏度时具有强铁磁性。在 723 摄氏度时，0.22 % C 溶解在铁氧体中，但在室温下，溶液中只剩下千分之一的 C。

渗碳体较脆，硬度较高（布氏硬度约为 800）。它在 210 摄氏度以下具有弱磁性，是电和热的不良导体。它有一个复杂的菱形格子。通常区分 (i) 初级 Fe3C，它从线 CD 处的液体结晶，(ii) 二级 Fe3C，它从线 ES 处的伽马溶液中沉淀出来，和 (iii) 三次 Fe3C，它从线 CD 处沉淀出来。 PQ 行的解。

石墨很软。它是电的不良导体，但可以很好地传递热量。石墨即使在3000℃到3500℃的温度下也不会熔化，它具有六方晶格，轴比c/a大于2。

奥氏体柔软（但比铁素体硬）且具有延展性。奥氏体的伸长率范围为 40% 至 50%。它的热导率和电导率比铁氧体低，并且是顺磁性的。奥氏体具有fcc晶格。

含 0 % – 0.02 % C 的钢的组织由铁素体和三级 Fe3C 组成。 C 含量的进一步增加导致出现新的结构成分，即铁素体和 Fe3C（珠光体）的共析体。珠光体首先作为铁素体晶粒之间的单独夹杂物出现，然后在 0.8% C 时占据整个体积。珠光体具有两相混合物的特征，通常具有层状结构。随着钢的C含量增加到高于0.8%的值，二次Fe3C与珠光体一起形成。次生Fe3C呈针状。 Fe3C的量随着C含量的增加而增加。在 2 % C 时，它占据显微镜视野的 18 %。当 C 含量超过 2% 时，出现低共熔混合物。在快速冷却的钢中，并非所有剩余相（铁素体或 Fe3C）在形成共析体之前都有时间沉淀。

含 3.6% C 的合金含有莱氏体（C 固溶在 γ-Fe 和 Fe3C 中的共晶混合物）。这些合金更适合与亚共晶白口铸铁分类。

临界（转变）温度

碳对固态铁的转变有显着影响。 Fe-C 平衡图中 GS 和 NL 线的位置 s 表明，相对于图 1 所示的对应物，C 含量的增加导致 A3 点降低和 A4 点升高。因此，C 延伸δ相的温度范围。

当形成共析体（珠光体）时，加热和冷却曲线停止。这被标记为 A1 点（Ac1 加热，Ar1 冷却）。这种现象发生在 0 .9 % C（Fe-C 图中的 S 点）。亚共析钢中铁素体的析出（穿过 GOS 线）在加热和冷却曲线中显示为拐点，由点 A3 表示。该点对应于纯铁中的 gamma 到 alpha 转换。在共析沉淀之前的 Fe3C 沉淀（穿过线 ES）在冷却曲线中被视为弱拐点，指定为点 Acm（加热时为 Ac,cm，冷却时为 Ar,cm）。 C的添加对磁转变温度（A2点）影响很小。因此，线 MO 对应于低 C 含量钢的磁转变。在含有较高量 C 的合金中，这种转变发生在 GOS 线上，这对应于铁素体析出的开始。如果C含量高于S点对应的C含量，则磁转变与温度A1一致。

渗碳体经历磁转变。无论 C 含量如何，转变发生在 210 ℃ ～ 220 ℃ 的温度下。它没有明显的滞后现象，纯 Fe 在 A2 点的磁转变也是如此。

钢的结构转变

当钢要硬化时，将其加热到高温以将整个组织转变为奥氏体相，该奥氏体相是在高温下稳定的Fe和C的单相组织。如果这种加热的钢缓慢冷却，奥氏体转变为珠光体，这是室温下的平衡相。珠光体结构是一种退火结构，相对较软，物理性能较低。如果加热的钢很快冷却，就会形成称为马氏体的坚硬而坚固的结构，它是溶解在铁中的 C 的亚稳态相。该相可以被回火以产生更不脆的较低硬度的组织。中间冷却速率会产生其他结构，例如贝氏体，尽管这种类型的结构仅在合金钢中大量生产。共析C钢主要产生马氏体或珠光体，取决于冷却速度。

奥氏体珠光体转变

由于奥氏体中溶解的 C 的存在，阻碍了奥氏体的 fcc 晶格向铁素体的 bcc 晶格的转变。奥氏体晶格有足够的空间容纳单元中心的 C 原子。铁氧体的 bcc 晶格没有这个空间。因此，在从奥氏体转变为铁素体时，C 的溶解度显着降低。在 β 到 α 的转变过程中，几乎所有的 C 从奥氏体晶格中析出。根据亚稳态 Fe-Fe3C 图，它以渗碳体的形式析出。这种转变可以通过三个相互关联的途径来定义，即（i）γ-Fe晶格转变为α-Fe晶格，（ii）C以Fe3C的形式沉淀，以及（iii）碳化物的凝聚。

在 A1 点的温度下，路线 (i) 和 (ii) 的转变几乎同时进行，形成铁素体和渗碳体的层状混合物。溶解的 C 原子随机分布在晶格中。因此，Fe3C 在富含 C 的区域中成核，而在碳含量很少（如果有的话）的区域中形成铁素体。这种C的重新分布是通过扩散发生的，并且取决于温度和时间。

当C含量低于0.8%的亚共析钢经受缓冷时，转变开始于晶界处铁素体的形成。该晶界充当铁素体结晶中心。碳被迫进入微晶。随着铁素体沉淀，铁素体形成所需的浓度在中心体积中达到。当过共析钢（C 大于 0.8%）经受缓冷时，在穿过 ES 线时，Fe3C 开始在晶界析出。这里晶界也是结晶位点。

随着温度降低，γ-Fe 和 α-Fe 晶格中的 C 扩散速率迅速降低，因为扩散系数取决于温度。在适当的冷却速度下，过冷度可以提高到不能形成珠光体的程度。

在低温范围内，转变机制和所形成结构的特征仅取决于发生转变的温度。考虑过冷度，区分三个转变温度范围，即（i）珠光体范围，（ii）中间范围和（iii）马氏体范围。从一种转变机制到另一种转变机制的连续转变可以在这些温度范围内发生。转变过程很大程度上取决于钢中 C 和其他元素的含量。它们可以以较快的机制开始，以较慢的机制结束。

在珠光体范围内，转变的特征是同时形成铁素体和碳化物的混合物。游离铁素体或碳化物可在奥氏体晶界析出。在这里，两相的形成和生长都受扩散过程的控制。 Fe和其他元素的扩散起着重要作用。随着温度的降低，组织细度提高，直到铁素体和碳化物的扩散结晶需要更长的时间。

珠光体是铁素体和碳化物板的机械混合物，在珠光体范围内转变形成。珠光体晶核的形成速度取决于奥氏体中碳化物的过饱和度，该过饱和度随着温度的降低而增加。该速率还取决于随温度降低的扩散速率。珠光体岛的生长主要取决于C和Fe原子的扩散速率。其他因素是（i）过饱和度和（ii）铁素体形成过程中的自由能优势。珠光体岛不仅通过新板块的形成而生长，而且还通过旧板块向各个方向进一步生长。硬质合金板比铁氧体板长得快。

珠光体的形成过程始于铁素体晶核的形成。铁素体和渗碳体板形核的多次交替以及两相板的分支导致形成平面平行和扇形珠光体板。珠光体核主要出现在具有晶体结构缺陷的晶格区域，如晶界、不溶性碳化物或非金属夹杂物。珠光体的一个非常显着的特征是板间距。钢的强度性能随着间距的减小而提高。

随着温度的降低，珠光体范围内的 Fe3C 和铁素体结晶中心的形成速度加快。结构细度越高，板间距越小。

影响钢性能的一个重要特征是珠光体团的尺寸。菌落尺寸的减小伴随着冲击强度的增加和脆性的降低。临界脆化温度取决于珠光体形态。从而在铁素体和渗碳体板断裂的情况下形成强度较高的珠光体，在铁素体内部形成高密度的位错。

珠光体更好的断裂强度是通过 Fe3C 颗粒的球化来实现的。通过珠光体的变形、随后的加热和保持在Ac1附近的温度可以促进球化。另一种提供珠光体相对高强度和延展性的方法在于珠光体转变期间的变形。这导致渗碳体形成多边形结构和球状化。铁素体-珠光体混合物的屈服应力（YS）取决于铁素体和珠光体的相加性质。

奥氏体转变

在亚共析钢和过共析钢中的奥氏体转变过程中，珠光体转变之前是过量相即铁素体和二次渗碳体的析出。结构自由过剩相的相对量取决于奥氏体过冷度。过量铁素体或 Fe3C 的量随着冷却速度的增加而减少。当过冷度足够时，可以避免形成作为独立结构成分的过量相。

当含有少量共析奥氏体的亚共析钢经受缓慢冷却时，共析铁素体在过量铁素体的晶粒上生长，并且共析Fe3C在晶界处作为结构自由中间层留下。在过共析钢中，共析钢也可能发生结构退化。由于在低于 A1 点（高于 700 摄氏度）的非常低的冷却下由于共析沉淀而形成的渗碳体沉积在二次渗碳体上。旁边注意到结构上游离铁素体的区域。这种伴随着相分离的共析转变被认为是异常的。在正常的共析转变中，铁素体和 Fe3C 以菌落的形式一起生长，两相有规律的交替。在异常转变的情况下，铁素体和 Fe3C 的粗混合物不具有独特的共析结构。在共析转变过程中，机制可以从异常变为正常。因此，通过快速冷却和相应显着的奥氏体过冷度，可以完全抑制异常转变。

在亚共析钢中铁素体过量的情况下，铁素体以两种形式存在，即 (i) 致密的等轴晶粒，和 (ii) 取向的 Widmanstätten 板。亚共析铁素体的致密析出物主要出现在奥氏体晶界，而维德曼板则在晶粒内部形成。 Widmanstätten 铁素体仅在碳含量低于 0.4% 且奥氏体晶粒相当粗大的钢中观察到。随着奥氏体晶粒尺寸的减小，等轴晶粒形式的铁素体份额增加。 Widmanstätten 铁素体形成于A3（50℃）至600℃至550℃的温度区间。随着钢中C含量的增加，Widmanstätten 铁素体在组织中的份额降低。

Widmanstätten 铁素体的形成应该是由于晶格的剪切 gamma-alpha 重排，伴随着原子的有序相互关联的运动。铁素体的等轴晶粒通过晶格的正常扩散重排和跨越γ/α边界的原子无序跃迁而生长。

用于强化钢的方法之一在于提供具有亚共析铁素体的结构，该亚共析铁素体含有分散的碳化物沉淀物。为了产生这样的组织，钢要加热到特殊碳化物溶解在奥氏体中，然后快速冷却，以防止碳化物在亚共析铁素体开始形成之前直接从奥氏体中析出。

马氏体相变

马氏体的转变是由于高温相的淬火（快速冷却）。 C钢中马氏体相变的主要特征如下。

• 马氏体转变是由于钢从高于 A1 的温度（例如在水中）快速冷却而发生的。由于快速冷却，抑制了奥氏体扩散沉淀为铁素体和碳化物的两相混合物。马氏体中 C 的浓度与奥氏体中的相匹配。马氏体转变发生时没有任何扩散。
• 奥氏体向马氏体的转变从马氏体起始温度 (Ms) 开始。 Ms 通常不依赖于冷却速度。马氏体是在一定的温度区间内形成的。具体温度由钢的C含量决定。
• 在温度区间 Ms-Mf（马氏体完成）内终止冷却会暂停马氏体的形成。该特征将马氏体转变与珠光体转变区分开来。在珠光体相变中，相变在 A1 点以下的恒定温度下持续到最后，在足够的等温保温时间下，最终结果是奥氏体完全消失。随着马氏体相变，留下一定量的残余奥氏体。
• 马氏体转变没有潜伏期。在温度Ms以下瞬间形成一定量的马氏体。
• 在低于 Ms 的冷却时，由于新板的快速形成，马氏体量迅速增加。最初形成的板块不会随着时间的推移而增长。
• 马氏体晶格相对于奥氏体晶格有规律地取向。晶格之间存在一定的取向关系。

Ms-温度表征了经过特定预处理的特定成分的钢。在给定的钢中，无论冷却速度如何，马氏体转变都在相同的温度下开始。该温度取决于钢的成分，并随着钢中 C 含量的增加而大大降低。部分C进入与奥氏体共存的碳化物。如果提高淬火温度，碳化物会溶解在奥氏体中。因此，奥氏体的C浓度增加，Ms点降低。

马氏体的形成被认为是奥氏体晶格重排的剪切机制。众所周知，相变的马氏体（剪切）机制是通过原子有序的相互关联运动到比原子间距更短的距离，并且原子不交换位置。初始相中的原子在马氏体相中保留其邻居。这是晶格剪切重排特有的主要特征。

晶格重排的这种性质提供了新旧相之间边界的连贯性。在马氏体和初始相之间的边界处晶格的相干性（弹性共轭）确保了即使在低温下也可以非常快速地向基体移动边界。原子协同移动到比原子间距更短的距离，从而导致马氏体晶体的生长。

随着马氏体晶体的生长，弹性应变在相干边界处累积。到达 YS 时，连贯性受到干扰。原子在马氏体晶体和起始基体之间的边界处变得无序。边界的滑动运动变得不可能。因此，通过马氏体机制的晶体生长终止，之后晶体只能通过扩散生长。但马氏体转变发生在低温下，扩散速率非常小。因此，相干性破坏后，几乎没有观察到马氏体晶体的生长。

马氏体机制引起的固溶体多晶相变的特点是不存在组分的扩散再分布。这里描述了高温相转变为低温相的马氏体机制所必需的条件。马氏体相变不可能在过冷度较小的情况下进行。这是因为在晶格无序重排的情况下，弹性变形仅由体积变化决定，而马氏体转变还取决于初始晶格和马氏体晶体的晶格的相干性。随着过冷度的增加，晶格的无序重排率增加，达到最大值，然后下降。为了得到Fe多晶相变的马氏体机制，钢要在伽马范围内强烈过热，然后快速冷却以抑制正常相变的发展。

在马氏体形成过程中，奥氏体的 fcc 晶格重排为马氏体的 bcc 四方晶格，类似于 α-Fe 的 bcc 晶格。奥氏体晶格通过贝恩变形转变为马氏体晶格，该变形包括沿 c 轴压缩奥氏体的四方晶胞并同时沿 a 轴增加尺寸。马氏体晶格的四方畸变程度c/a直接随着马氏体的C浓度而增大。马氏体晶格在室温下保持四方性。建立了初始相和马氏体相的取向关系。

关于马氏体形核的性质有许多假设。他们中的许多人主张在起始基质的特殊缺陷位点进行异质成核。

马氏体在形态上分为两种基本类型。这些是板状马氏体和块状马氏体。它们在形状、晶体的相互排列、亚结构和习性平面上都有所不同。板状（针状）马氏体更常见于高C钢中。马氏体晶体呈薄透镜状板状。首先出现的板通过整个单元，将其分成单独的部分。但它们不能穿过基体晶界。因此，板尺寸受到奥氏体晶粒尺寸的限制。在奥氏体截面中形成新的马氏体板。此处板尺寸仅限于截面尺寸。如果奥氏体晶粒小，马氏体板很细，以至于在显微截面试样中看不到马氏体的针状组织。这种马氏体称为无组织马氏体，是最理想的。

在低碳钢和中碳钢中可以观察到块状（板条）马氏体。这种类型的马氏体晶体被成形为具有大致相同取向的互连板。块状马氏体板块以低角度边界分隔。

贝氏体的转变

贝氏体转变介于珠光体转变和马氏体转变之间。贝氏体的相变动力学及其形成的组织呈现出扩散珠光体相变和少扩散马氏体相变的特征。由于这种转变，形成了铁素体和碳化物的混合物。这种混合物称为贝氏体。贝氏体转变机制包括晶格的γ到α重排、C的重新分布和碳化物的析出。

这里解释了贝氏体转变与其珠光体和马氏体对应物的接近程度。基本成分 Fe 的原子扩散运动在贝氏体转变范围内几乎完全被抑制。然后由于珠光体析出的抑制，铁素体的 γ 到 α 形成是困难的。然而，C 扩散相当活跃并导致碳化物沉淀。在中间范围内，伽马相晶体是通过像马氏体板一样的连贯生长形成的。但α相板的形成是缓慢而不是瞬间形成的。

这是因为在中间温度范围内，α 相只能从 C 耗尽的 γ 相中沉淀出来。 Thus the growth rate of the alpha phase crystals depends on the C diffusive removal rate. In this case, the martensite start point Ms in austenite rises and the martensite gamma to alpha transformation takes place at temperatures above the temperature Ms typical of the steel with a given composition.

At the instant of martensite transformation, the C concentration remains unchanged. Only the crystal lattice is altered and a supersaturated a solution is formed. Carbide precipitates after gamma to alpha transformation.

There is a difference between upper and lower bainite, which are formed in the upper and lower parts of the intermediate temperature range. The conventional boundary between the bainite is close to 350 deg C. Upper bainite has a feathery structure, whereas lower bainite shows an acicular morphology, which is close to that of martensite. The difference in the structures of upper and lower bainite is due to the difference in the mobility of C in the upper and lower parts of the bainite temperature range.

The alpha phase substructure of upper bainite resembles the substructure of massive martensite in low C steel, while the alpha phase structure of lower bainite approaches the structure of martensite in high C steels. In upper bainite, carbide particles can precipitate both at lath boundaries and inside laths. This fact suggests that here carbides precipitate directly from austenite. In lower bainite, carbide is found inside the alpha phase. This is since carbide is formed during precipitation of a supersaturated solid solution of C in the alpha phase. Both upper and lower bainite shows a high density of dislocations inside the alpha phase. Fe3C is the carbide phase in upper bainite and epsilon carbide (Fe2C) in lower bainite. As the holding time is increased, Fe2C turns into cementite. The dimensions of austenite grain have no effect on the kinetics of martensite transformation.

Tempering

The processes which take place during tempering are precipitation and recrystallization of martensite. Quenched steel has a metastable structure. If subjected to heating, the structure becomes closer to equilibrium. The nature of the processes which occur during tempering is determined by three major characteristics of quenched steel namely (i) strong super saturation of the martensite solid solution, (ii) high density of crystal lattice defects (dislocations, low angle and large angle boundaries, and twin interlayers  etc.), and (iii) presence of retained austenite.

The main process taking place during tempering of steel is the precipitation of martensite accompanied by formation of carbides. Depending on the temperature and duration of tempering, the martensite precipitation can involve three stages namely (i) pre-precipitation, (ii) precipitation of intermediate metastable carbides, and (iii) precipitation and coagulation of cementite. Retained austenite can precipitate simultaneously. Since there is a high density of dislocations in martensite, hence its substructure is similar to the substructure of steel which is work hardened. Hence, polygonization and recrystallization can develop during tempering.

When C steel is tempered, super-saturation of the gamma solution in austenite increases with an increase in the C content of steel. This leads to lowering of the Ms-temperature and transition from massive martensite to plate martensite. The amount of retained austenite also increases.

The segregation of C represents the first structural changes which take place during tempering of C steel. The segregated C can nucleate heterogeneously at lattice defects or homogeneously in the matrix. The heterogeneous nucleation of the segregated C takes place either during quenching or immediately after it.

Flat homogeneous clusters of C atoms not connected with lattice defects are formed at tempering temperatures of less than 100 deg C. This is due to the considerable displacements of Fe atoms and the appearance of elastic distortions. As the tempering temperature is increased, the clusters become larger and their composition is close to Fe4C. This process is dependent on the C diffusion. Metastable Fe2C is formed above 100 deg C. It has a hexagonal lattice and appears directly from C clusters when the C concentration is increased. Metastable Fe2C can also precipitate directly from the alpha solution. Fe2C precipitates as very fine (10 nm to 100 nm) plates or rods at low temperatures. With an increase in tempering temperature or time, Fe2C particles become coarser and precipitate in steels containing a minimum of 0.2 % C. In steels with a high Ms-temperature, partial precipitation of martensite is associated by the deposition of excess carbide and is obtained during quench cooling in the martensite range. Hence self-tempering of these steels occurs during their quenching.

Cementite is formed at a temperature higher than 250 deg C. Two known mechanisms of Fe3C nucleation are (i) precipitation directly from a supersaturated alpha solid solution and growth of Fe3C particles at the expense of the dissolution of less stable carbides, and (ii) appearance of Fe3C as a result of transformation of the intermediate carbide lattice to the Fe3C lattice.

In the final stage of the carbide formation during tempering, coagulation and spheroidization of carbide take place. These happen intensively starting from 350 deg C to 400 deg C. At temperatures higher than 600 deg C, all Fe3C particles have a spherical shape and undergo coagulation only.

A substantial part of the tempering process is devoted to the precipitation of retained austenite accompanied by deposition of carbides. Precipitation occurs over the temperature range of 200 deg C to 300 deg C. During tempering, retained austenite transforms into lower bainite.

A decrease in the C concentration of the alpha phase during carbide formation results into changes in the phase structure. Martensite precipitation is conventionally divided into two stages. The first stage of precipitation is achieved below 150 deg C when the mobility of C atoms is sufficient for the formation of carbide plates. But, it is insufficient for the carbide plates to grow by diffusion of C from the areas of non-precipitated martensite with a high C concentration. This results in a non-uniform content of C in different areas of the martensite and hence inhomogeneity of martensite results with respect to its tetragonality. In areas with precipitated carbide, tetragonality is lower than in non-precipitated areas. Two solid solutions with different C concentrations coexist. For this reason the precipitation is referred to as a two -phase precipitation. The two phase precipitation of martensite results from the deposition of new carbide particles in areas containing martensite with the initial C concentration. Carbide particles do not grow at this stage.

At the second stage of martensite precipitation (150 deg C to300 deg C the alpha solution is depleted of C owing to diffusive growth of carbide particles, but the process proceeds very slowly. Hence, the precipitation kinetics is due to the rapid depletion of the alpha solution in carbon.  Subsequently, depletion of the solid solution in C stops. At 300 deg C around 0.1 % C is left in the alpha solution. Above this temperature, no difference between the lattice of the alpha solution and that of the alpha-Fe is detected. Below 300 deg C the degree of tetragonality is still measurable. Above 400 deg C the alpha solution becomes completely free of excess C and transformation of martensite to ferrite is finished.

Plates (needles) of quench martensite have a high density of dislocations which is comparable to the density of the deformed steel. However, recrystallization centres and their progress to recrystallized grains are not observed. This is since carbide particles pin dislocations and large angle boundaries. It is only above 600 deg C, when the density of the particles decreases owing to the coagulation, that the recrystallization growth of grains takes place at the expense of migration of large angle boundaries. With this the morphological structures of lath martensite disappear. These processes are hampered in high C steels as compared to low C steels, since the density of carbides is greater in high C steels. The acicular structure is retained up to the tempering temperature of around 650 deg C.

The structural changes which occur during tempering cause alteration of steel properties. These changes depend on the tempering temperature and time. Hardness decreases as the tempering temperature is increased.

Kinetics of transformation of austenite

The kinetics of transformation of austenite is described below.

Isothermal transformation diagrams

It is important to follow the process at a constant temperature for the understanding of the kinetics of the transformation to austenite. For this purpose, isothermal transformation (IT) diagram is usually made which illustrates the isothermal process of austenite precipitation. In IT diagram (Fig 2), the transformation time is in the X-axis shown on the logarithmic scale and the temperature is plotted on the Y-axis. From this diagram, the incubation period (left hand curve) can be determined and also the time required for completion of the process (right hand curve). The instant, steel passes the points A3 and A1 during quenching, is usually taken as the zero time reference.

The time required to achieve the temperature of the quenching medium is frequently neglected. The start and finish of the transformation are difficult to determine from the transformation curve behaviour at the initial and final sections of the curve. Hence, the lines of the IT diagram generally correspond to a certain final volume which has undergone transformation, e.g., 2 % and 98 % for the transformation start and finish, respectively. The volume value is usually not shown in the IT diagram.

Fig 2 Isothermal transformation diagram

In addition to the curves stated above, IT diagram frequently contains intermediate curves corresponding to certain values of the transformed volume, say 20 %, 50 %, or 80 %. A decrease in the transformation rate causes displacement of the transformation start and finish curves to the right, i.e., toward greater duration. This phenomenon can be seen if the quenching heating temperature increases as a result of a decrease in the number of inclusions and growth of austenite grains. An increase in the transformation rate leads to displacement of the curves to the left. This phenomenon can be accounted for (i) by a decrease in the quenching heating temperature, (ii) the presence of carbides or inclusions, and (iii) refinement of the austenite grain. For a specified sample of steel the temperature which corresponds to a maxi mum transformation rate (the nose of the sigmoid curve) does not, as a rule, change significantly.

Continuous cooling transformation diagrams

Continuous cooling transformation (CCT) diagrams consider the transformation kinetics of eutectoid steel. The major transformation which takes place during annealing cooling of steel is a eutectoid precipitation of austenite into a mixture of ferrite and carbide. The eutectoid transformation kinetics is given by IT diagrams of austenite at a temperature of 727 deg C. The structure attained after tempering below 300 deg C is called tempered martensite. An acicular structure is seen after tempering at 300 deg C to 450 deg C. Tempering over the temperature interval of 450 deg C to 600 deg C shows a distinct dot structure. Austenite is in a thermodynamically stable equilibrium with the ferrite-Fe3C mixture. Stability of undercooled austenite is defined by a period of time during which the appearance of precipitation products in the diagram cannot be registered by conventional methods. The degree of austenite undercooling is the main factor which determines the steel microstructure. The necessary degree of undercooling is provided by either continuous cooling or isothermal treatment.

As seen earlier, in hypo-eutectoid steels the formation of pearlite is preceded by precipitation of hypo-eutectoid ferrite. With a decrease in the transformation temperature and an increase in the degree of undercooling, precipitation of hypo-eutectoid ferrite is suppressed. The amount of pearlite increases and the C content becomes less than that in pearlite of the eutectoid steel. In the region of the maximum transformation rate, the two curves merge. Thus, a purely pearlitic structure is formed in steel with 0.4 % C. In steels containing higher amounts of C, the precipitation of ferrite cannot be suppressed even if the C content decreases. Ferrite precipitation precedes the formation of pearlite even at a maximum transformation rate, but the amount of ferrite is less than that is formed at smaller undercooling.

These propositions are valid for the precipitation of cementite in hyper-eutectoid steels, but it can be suppressed even at relatively small undercooling. In this case, the C content of pearlite becomes higher than that in the eutectoid steel. As a result of suppression of the hypo-eutectoid ferrite precipitation under continuous cooling from the region of the gamma solid solution, the point Ar3 lowers much faster than the point Ar1 as the cooling rate is increased. With a certain cooling rate, both points merge into one point, which corresponds to the formation of a fine plate structure of the pearlite type free of ferrite.

Under continuous cooling the transformation process can also be visualized as diagram in temperature-time coordinates. Therefore the behaviour of cooling curves is to be analyzed to find characteristics of the transformation processes. In this diagram, the ferrite and pearlite start lines are shifted toward longer periods of time compared to the IT diagram. This is due to an increase in the temperature interval necessary for preparing the transformation processes in the austenite lattice. As a result, only part of the incubation period, which is needed for the IT to start, is effective. In this case, the incubation period is the mean of the effective lengths of time corresponding to different periods of time in the given range. This proposition can be used to calculate the behaviour of the transformation start line in the pearlite range from the IT diagram. The reverse calculation is also possible.

Similar to the pearlite range, in the bainite temperature range, the precipitation of undercooled austenite starts after a certain incubation period. Resemblance of the bainite and pearlite transformation kinetics consists not only in the presence of an incubation period but also in the character of the volume increase during isothermal soaking which is the fraction of the transformed volume of austenite increases first with acceleration and then with deceleration. At the same time, as in the case of the martensite transformation, retained austenite does not disappear completely during the bainite transformation. Every point in the bainite finish curve corresponds to certain amount of retained austenite. Similar to the pearlite transformation, the bainite transformation can take place both during isothermal soaking and under continuous cooling. Austenite which has not been transformed over the bainite range turns partially into martensite when the steel is cooled to room temperature. Since the austenite is inhomogeneous with respect to the C content after the bainite transformation, martensite is formed predominantly in C enriched regions.

For the high alloy steel, IT curves can be separated by a temperature interval in which undercooled austenite is highly stable. In this interval, pearlite precipitation does not take place for many hours, while undercooling is inadequate for the bainite transformation. In C steel, the bainite transformation proceeds concurrently with the pearlite transformation. Products of the pearlite transformation dominate at higher temperatures, and those of the bainite transformation at lower temperatures.

During the transformations of austenite on cooling in the martensite range, martensite component in the steel structure appears when the cooling rate achieves a certain value. The minimum cooling rate at which the martensite component is formed is called the lower critical rate of cooling. The rate at which transformations by the pearlite and bainite mechanisms are suppressed completely is referred to as the upper critical rate of cooling (quenching). If the conditions of austenite formation (austenitization temperature and the holding time at this temperature) and the cooling conditions (cooling rate exceeds the upper critical rate) are constant, the location of the martensite start point Ms depends only on the contents of C and alloying elements in the steel.

If the cooling rate is high, the formation rate of separate needles of martensite is also high, and transformation of austenite to martensite begins on reaching Ms-temperature. It continues on subsequent cooling to lower temperatures. As the temperature of the quenching medium is lowered, the amount of formed martensite increases first quickly and then slowly. With an increase in the quenching heating temperature (austenitization temperature), the transformation also shifts toward lower temperatures as more of the alloying elements are taken into solution. A certain amount of martensite can be formed during isothermal holding, but it is not high in C steels. Retained austenite is stabilized during isothermal holding. As a result, more martensite is formed during subsequent cooling. Formation of martensite stops at the point Mf. There is a relationship between some factors which influence the stabilization of martensite. The effect of stabilization increases with the amount of martensite in the structure or, the amount of martensite being equal, with temperature.

There is a close link between the CCT and IT diagrams. When resolving practical issues involved in heat treatment of steel, it is sometimes necessary to know how the continuous cooling rate affects the structure formed as a result of austenite transformation. For this, there have been efforts to establish the relationship between the transformation kinetics of austenite under isothermal conditions and under continuous cooling conditions. The efforts have started from the concept of additivity of the transformation processes at different temperatures. It has been presumed that holding of undercooled austenite at a preset temperature is part of the incubation period. However, it has been found, that calculated and experimental data coincide satisfactorily only if the pearlite transformation is continuous.

If the pearlite transformation is preceded by precipitation of eutectoid pearlite or the pearlite and bainite transformations occur concurrently, calculated data are at a discrepancy with the experimental data. It has been found that the discrepancy is due to the factors namely (i) holding of austenite during the time accounting for fractions of the incubation period causes acceleration of the subsequent intermediate transformation at the expense of preparatory processes, (ii) precipitation of hypo-eutectoid ferrite alters the austenite composition which delays the subsequent intermediate transformation, (iii) partial transformation of austenite over the intermediate range reduces the rate of the said trans formation at lower temperatures and facilitates an increase in retained austenite which is due to a redistribution of C and enrichment of the non-transformed part of austenite in carbon, and (iv) a change in the cooling rate over the martensite range affects stabilization of austenite in different ways.

For the above reason, special methods of constructing thermo-kinetic transformation diagrams of austenite subject to continuous cooling have been elaborated for non-eutectoid steels. From these diagrams it is possible to determine the critical rate of quenching cooling or continuous cooling which is necessary to complete a particular stage of austenite precipitation.

It has been seen that the CCT diagram is a function of the bar diameter. When steel is subjected to martensitic hardening, it is required to be cooled from the quenching temperature so that on undercooling to a temperature below the Ms point austenite has no time to precipitate and form a ferrite-carbide mixture. For achieving this, the cooling rate is to be less than the critical value. The critical cooling rate is the minimum rate at which austenite does not precipitate to a ferrite-carbide mixture. Of course, the cooling rate of steel products is non-uniform over their cross section. It can be higher than the critical rate on the surface and lower than the critical rate at the centre.

The critical cooling rate at different points of a product can be directly determined from an IT diagram. In the first approximation, it is given by the slope of the tangent to the C curve which denotes the austenite precipitation onset. This method gives a value which is around 1.5 times the true critical rate. The cooling rate can be determined more accurately if thermo-kinetic diagrams are used. Intercepts of the cooling curves with the lines of the thermo-kinetic diagrams show the start and finish temperatures of the corresponding transformation.

From the transformation diagram, it is possible to determine, for example, the rate which provides 40 % martensite in the structure or the rates at which the entire transformation occurs in the pearlite range, i.e., hardening is omitted altogether. Because the data on the critical hardening rate depend on cooling time and is to be associated with a specific temperature (at which direct measurements of the hardening rate are practically impossible), it is proper to specify the cooling time for a specific interval of temperature, for example, from the point A3 to 500 deg C. Point A3 in the diagram is the time reference. Then it is possible to directly determine the critical cooling time K (Km for fully martensitic hardening, Kf for initial appearance of ferrite, and Kp for full transformation in the pearlite range).

Since the cooling time and the progress of the subsequent cooling of the sample during end-face hardening are known, the outcome of hardening can be determined from the transformation diagram. It is to be remembered that a transformation diagram is valid only for particular conditions of melting and homogenization. Deviations in the composition or grain dimensions cause changes in the trend of thermodynamic curves. This is explained by the fact that an increase in the homogenization temperature and time and, consequently, enlargement of the grains enhance the stability of austenite. Conversely, refinement of grains lowers the critical cooling rate, since stability of austenite decreases with an increase in the extent of grain boundaries.

Hardenability

The depth of the hardened zone is termed hardenability. This is one of the most important characteristics of steel. Since the cooling rate is non-uniform along the cross section of a sample, austenite can pass into martensite in surface layers only, while at the centre of the sample austenite undergoes the pearlite transformation. In the first place, hardenability depends on the critical cooling rate. An examination of the temperature curves plotted for different areas of the sample shows that the cooling rate of the core of a large diameter product is lower than the critical value and hence the core is not martensitically hardened. Martensite is present in the surface layer only.

After hardening treatment, a bulky part with a large cross section can show the entire range of structures such as a smooth transition from martensite near the surface through troostite-martensite and troostite to pearlite at the centre. The geometry of samples can influence the character of the cooling curves. However, given the same surface-to-volume ratio, the curves coincide in general. The highest changes in the cooling rate are experienced by the diameter of samples.

Considering the above, for achieving a through hardening of bulky products or full martensitic hardening to the core of a product, it is essential to provide the critical hardening rate along the entire cross section of the product. IT and CCT diagrams can be used to determine this rate. The diagrams are usually plotted for different grades of steel, taking into account the progress of cooling in different sections and in different hardening media.

The hardenability of steels depends on the steel composition, specifically on the C content. In the steel hardenability diagrams, the hardenability of each grade of steel is normally presented as a hardenability band. These diagrams have been plotted for almost all existing grades of steel. They show how to achieve hardening of a product made of particular steel.

Hardenability of steel is also categorized by IT curves. The more the curve is shifted to the right along the X-axis, the greater is the hardenability of the steel. This is explained by the fact that the rightward shift of the IT curve is due to better stability of austenite.

An improvement in the stability of undercooled austenite and hence an increase in the critical hardening rate lead to a greater depth of hardening. Then hardenability depends on all the factors which improve the stability of undercooled austenite. As an example, the stability of austenite can be raised by alloying steel with chromium and tungsten. These elements lower the austenite precipitation rate and can make steel an air-hardening one. Steel with a normal content of impurities is hardened to strength ten times that of a pure Fe-C alloy.

Elevation of the hardening temperature favours an increase in the hardening depth due to the homogenization of austenite and enlargement of austenite grains. Refinement of grains impairs hardenability as grain boundaries affect the stability of austenite. The hardening depth also depends on the hardening medium used. The greater is the intensity of cooling, the greater is the depth of hardening. Besides, the hardening depth depends on the cross-sectional diameter of the products. The critical diameter is that of the greatest cross section which lends itself to through hardening in a given hardening medium. The critical diameter is different for different hardening media and characterizes the hardenability provided by a particular method only.

Hardenability has an effect on the mechanical properties of steel. In the case of through hardening, the properties do not differ along the cross section of a product. Otherwise they decrease from the surface to the centre.  The analysis of the influence of hardenability on the properties of steels which have been tempered after hardening shows that a high temperature favours equalization of hardness along the cross section. However, the structure of weakly hardenable steels remains inhomogeneous. This is due to a grain structure appearing on the surface, where martensite is formed during quenching, while a lamellar structure remains at the centre. A grain structure is present along the entire cross section of through-hardening steel. This determines the character of changes in the properties of steels with different hardenability. The properties which are independent of the Fe3C form (YS, specific elongation, impact strength) differ.

The properties of tempered steels (fracture stress, YS, impact strength, reduction in area) are impaired if ferrite precipitates during quenching. The mechanical properties of a product depend on its cross-sectional area. To obtain the best mechanical properties in the tempered state, a grain structure is required to be provided along the entire cross section; i.e., through hardenability is to be ensured in the quenched state.

Grain size

It is necessary to know the material structure while analyzing any processes or properties associated with grain boundaries. Most of the steel materials have polycrystalline structure and they comprise a set of grains separated by boundaries. The grain boundary is one of the basic structural elements in polycrystalline steel materials. The grain boundary represents an interface between two differently oriented crystals. This is the region of crystal imperfection. It is capable of moving and adsorbing impurities. The boundary has a high diffusive permeability.

In polycrystalline steel materials, the boundaries determine the kinetics of many processes. For example, movement of grain boundaries controls the process of recrystallization. A high diffusive permeability of grain boundaries determines the kinetics of diffusion-dependent processes at moderate temperatures.  Embrittlement of steel material is connected with enrichment of grain boundaries in impurities.

Grain boundaries are normally divided into two large groups namely (i) low angle boundaries, and (ii) large angle boundaries. Low angle boundaries are sub-grain boundaries with an angle of less than 10 degrees. They represent networks or walls of dislocations. The structure of large angle boundaries is much more complicated. The progress in understanding the structure of grain boundaries is connected with elaboration of the models describing the observed microscopic properties of the boundaries.

Grain size determination

The size of the grain that is formed under a given treatment is determined from micro-sections after their etching. For C and alloyed steels the reagent used is 1ml to 5 ml HNO3 +100 ml ethyl or methyl alcohol. Austenitic steel is etched in a copper sulphate-chloride solution containing 10 grams copper sulphate, 50 ml hydrochloric acid, and 50 ml water. When C and low alloy steels are etched, the reagents turn pearlite dark and make visible the ferrite grain boundaries, the martensite structure, and tempering products. The etching rate rises with the amount of nitric acid. The etching time is from several seconds to a minute. Etching of austenitic steel reveals the austenite structure and the austenite grain boundaries.

Carburization is also used to establish the austenite grain boundaries. In this case, samples are heated to 930 deg C in a carburizing medium (e.g., a mixture of 40 % BaCO3 and 60 % charcoal), cooled, and etched.

In addition, an oxidation method is used according to which micro-sections are heated in vacuum to a temperature 20 deg C to 30 deg C higher than the quenching temperature and are soaked for 3 hours. Subsequently air is fed to the furnace for 30seconds to 60 seconds, and the samples are cooled in water. Before quenching it is desired to heat samples in borax melt at 930 deg C to 950 deg C for 30 seconds to 40 seconds and then cool them in water. After these treatments micro-sections are polished and etched in a 15 % solution of hydrochloric acid in ethyl alcohol. Grain boundaries are seen as the oxide network.

Apart from this, use is made of the method of etching austenite grain boundaries, the method of the network of ferrite (for steels with a C content of up to 0.6 %) or Fe3C (for hypereutectoid steels), and the method of the pearlite network for steels which are closer in composition to eutectoid steels.

The grain size is determined by comparing the observed microstructure at a 100x magnification with standard scales (the scales are elaborated so that at a magnification of 100x the grain number N corresponds to the formula ‘n =8 X 2 to the power n’, with n the number of grains per sq mm of the micro-section area) or by counting the number of grains per unit area of the micro-section, or by calculating the mean nominal diameter of the grains or their number per cubic millimeter.

The austenite grain boundary structure which is produced on heating above the critical points is important since the austenite transformation products formed during cooling (martensite and pearlite etc.) appear inside austenite crystals. A coarse austenite grain determines a coarse plate structure of martensite during quenching or a coarse cellular network of ferrite (cementite) precipitates at the boundary of the initial austenite grains during annealing or normalization. The pearlite structure is also the coarser and the larger is the pearlite grain.

As is known, a coarse grain structure of steel (ferrite-pearlite, martensite, etc.) is characterized by lower mechanical properties. For this reason a fine-grain structure of steel is desirable in practice.

Grain size refinement

It is possible to refine a coarse-grained structure and this is widely used in the heat treatment of steel. The grain refinement, which takes place on heating steels above the Ac3 temperature, is related to a transition to the austenite state through nucleation of numerous centres of the austenite phase. Development of these centres leads to formation of a relatively fine grained structure. Above Ac3 temperature, the cross sectional size of the grain is 10 mm -30 mm. Initially the grain size is independent of the grain of the starting structure. It can be very fine irrespective of whether the starting structure of the steel is fine or coarse. A fine grain structure of the restored austenite provides a fine grain structure of cooled steel irrespective of the structural components (pearlite, bainite, or martensite) which are formed. This is due to the fact that all the transformation products nucleate within each separate grain of austenite.

Excess phases (ferrite in hypo-eutectoid steel and Fe3C in hyper-eutectoid steel) precipitate at boundaries of small austenite grains, and the pearlite transformation is accompanied by the appearance of smaller pearlite colonies. Fine austenite grains determine the formation of fine-needle martensite. This underlies the grain refinement effect which is associated with heating above Ac3 temperature. Heating the steel above Ac3 temperature during full annealing, normalization, or quenching is followed by recrystallization. With an initially coarse grain structure, recrystallization results in refinement of grains at a heating temperature corresponding to Ac3 temperature.

If the heating temperature is much higher than Ac3 temperature, then the grain is enlarged again, and the expected correction of the structure during the gamma to alpha transformation does not take place. Refinement of crystallites is especially pronounced when transformation to the austenite state starts in many centres inside the initial structure. The formed centres are to have a random orientation, which is not connected with the orientation of the alpha phase in the initial structure. Normally such centres are sufficiently large in number so that the grain size does not exceed 15 mm to 30 mm. During pearlite precipitation of austenite, breaking of an austenite grain into pearlite colonies, each of which can be considered an independent grain, also represents refinement of steel.

Strengthening mechanism in steel

There are four strengthening mechanisms in steel namely (i) solid solution strengthening, (ii) grain size refinement, (iii) dispersion strengthening, and (iv) work hardening.

Solid solution strengthening is a phenomenon which occurs when the number of impurity atoms in the lattice of the basic element is so small that they are incapable of forming both stable and metastable precipitation phases under any heat treatment conditions. However the impurity atoms favour improvement of the mechanical properties. The presence of impurity atoms in the matrix lattice leads to distortion of the lattice because of the difference in size between the atomic radii of the impurity and the basic component. This in turn leads to the appearance of elastic deformation fields, which retard movement of dislocations in slip planes under the action of applied stresses. In addition, the impurity atoms can obstruct movement of dislocations by forming impurity atmospheres around them. Both of the above factors play a leading role in solid solution strengthening.

Carbon which is statistically uniformly distributed in the lattice of the alpha iron has an influence on the structure and properties of alpha iron. Solubility of C in alpha iron is much lower than in the gamma iron. It forms interstitial solid solutions with both types of irons. However, whereas the gamma iron lattice has sufficiently large pores for implantation of C atoms, the cubic lattice of the alpha iron suffers. Upon introduction of C atoms, a tetragonal distortion takes place which is similar to the one of the martensite lattice except that in the former case the distortion is much smaller. In addition, inserting of C atoms causes the entire lattice of the alpha iron to somewhat expand. Hence, C affects the properties of the alpha phase. Actually, there is a dependence of the YS on the C concentration in the solid alpha solution. The influence which C exerts on plastic deformation resistance of the alpha phase is due to its strong interaction with dislocations as well as pinning of the dislocations and elastic deformations arising as a result of the tetragonal distortion of the alpha phase lattice after insertion of C atoms.

The presence of C in lattices of different structural components formed during thermal treatment of steel also leads to changes in their mechanical properties. As an example, the location of inserted C atoms primarily in one of the sub-lattices of interstitial sites during the martensite formation brings about additional tetragonal distortions of the martensite crystal lattice. This enhances plastic deformation resistance owing to the interaction between the stress fields around C atoms and those at dislocations. The influence of C dissolved in the alpha phase on the mechanical properties of steel is also witnessed in the case of the ferrite – pearlite transformation. The dissolution of part of the C in the alpha phase suggests that the solid solution strengthening of the phase is one of the factors providing the high strength properties of intermediate transformation products.

Grain size refinement of steel has a strengthening effect on steel. Impact strength is especially sensitive to the austenite grain size, and it decreases with grain enlargement. A decrease in the dimensions of pearlite colonies inside the initial austenite grain also favours a rise in impact strength.

Although the grain size has a considerable effect on impact strength, its influence is small if any on the individual mechanical properties such as hardness, fracture stress, YS, and specific elongation. Only the actual grain size affects steel properties, the inherited size has no effect. However, the technological process of heat treatment is determined by the inherited grain.

In the steels, precipitation of supersaturated solid solutions formed during quenching is followed by precipitation of disperse particles enriched in atoms of the alloying components. The strength (hardness) of the steels increases with the precipitation of these particles. The increment in the value of these characteristics increases as the dispersion and volume fraction of the particles increase. This phenomenon has been referred to as dispersion strengthening.

Precipitation of supersaturated solid solutions takes place during the heating (aging) of quenched steels. The strengthening is due to an increase in resistance to the movement of dislocations in a crystal when obstacles (barriers) of any type are formed. In aging steels, dislocations meet regions which retard their movement. The character of interaction between moving dislocations and precipitates of the second phase can be different depending on the phase morphology and structure. The total effect of aging on the strength properties of steels is determined by (i) the strength of the precipitates formed, (ii) the volume fraction of precipitates, (iii) the degree of precipitate dispersion, (iv) morphology, structure, and type of binding with the matrix, and (v) temperature.

When a solid solution of C in alpha Fe is cooled below A1 temperature, C precipitates as Fe3C with lowering of the C solubility and a decrease in temperature. This process takes place under sufficiently slow cooling, which is accompanied by diffusion processes, leading to the formation of cementite. In the case of abrupt cooling (water quenching) C has no time to precipitate. A super-saturated alpha solid solution appears. During subsequent storage at room temperature (natural aging) C tends to precipitate from the solid solution. Carbon enriched regions appear primarily in defective sections of the matrix. Precipitation of C from a supersaturated solid solution during natural aging results in an improvement of its strength properties and hardness. However, plastic properties such as reduction in area, specific elongation, and impact strength are deteriorated and the phenomenon of dispersion strengthening is seen.

As the heating temperature is increased (artificial aging), dispersion strengthening accelerates. This is due to the intensification of diffusion processes with an increase in temperature. The total process of C precipitation from the super-saturated solid solution in alpha Fe comprises several successive processes. Mechanical properties and hardness are not sensitive to structural changes which take place during the aging of the steels. Sharp changes in properties indicate alterations in the structural state of the steel.

A maximum change in mechanical properties during precipitation is achieved only if excess crystals in a highly disperse state precipitate. Subsequent coagulation of the crystals leads to degradation of the properties.

The influence of different solubilities of C in alpha Fe on the properties of the steel (dispersion strengthening) during low temperature aging is prominent in low C steels. In steels containing C higher than 0.4 %, the above effects are not noticed due to the influence of Fe3C particles formed during the pearlite transformation. Besides, nucleation of the precipitating phase can be inhibited owing to migration of C to the Fe3C-ferrite interfaces. As a result, the amount of C concentration at lattice defects decreases.

Cold plastic deformation greatly accelerates precipitation of a supersaturated solid solution. This is due to an increase in the density of dislocations, which are preferable sites of heterogeneous nucleation of precipitates as well as to an increase in the concentration of vacancies, which facilitates the diffusion of C to clusters. Mechanical properties change during aging after cold working in the same way as after quenching, that is, the YS, the fracture stress, and hardness are altered. With an increase in aging time, specific elongation and reduction in area decrease and the tendency to brittle fracture is enhanced. The rate of change is higher than in quenched steel. Also, the nature of the changes is different. Whereas in the case of aging after quenching, hardness reaches a maximum and then drops, after cold working hardness does not decrease with the aging time. As the aging temperature is raised, the maximum hardness of quenched steel lowers, while after cold working hardness is independent of the aging temperature. This is explained by the fact that a considerable amount of C is concentrated near dislocations. Few, if any, clusters nucleate in the matrix homogeneously. Consequently, clusters cannot grow at the expense of other clusters, i.e., they cannot coagulate.

An important method used to strengthen steels is deformation strengthening. Strengthening achieved with crystal deformation can be judged from the shape of stress-strain curves. The actual shape of these curves largely depends on the crystal lattice type of the metal, its purity, and thermal treatment.

In the case of cubic lattice steels, strengthening curves are parabolic, whereas for hexagonal lattice metals a nearly linear dependence is observed between the stress and the strain. This fact suggests that plastic deformation strengthening is determined mainly by the interaction of dislocations and is associated with the structural changes which retard the movement of dislocations. Metals with a hexagonal lattice are less prone to deformation strengthening than cubic lattice steels because the hexagonal lattice has fewer easy slip systems. In cubic lattice steels, the slip proceeds in several intersecting planes and directions.

There are three stages during the work hardening. The first stage is due to the easy slip. It depends on the orientation of the crystal relative to external forces and on the presence of impurities. This stage is characterized by a linear dependence of strain stresses on the strain at a small work hardening rate. Dislocations usually slip in primary systems.

In the second stage the work hardening rate is much higher than the first stage. Dislocations move in intersecting slip planes and, on colliding, form additional obstacles to their movement. This state is most extensive in the stress-strain curve. The ratio between the work hardening rate and the shear modulus (or any other elastic constant) is almost independent of the applied stress and temperature. It depends little on the crystal orientation and presence of impurities.

In the third stage changes are possible in the distribution of dislocations. They can either get around obstacles which retard their movement at the second stage or interact with dislocations. As a result, the work hardening rate is lower compared to which is observed during the second stage. At this stage, a partial relaxation of stresses can occur owing to the appearance of the secondary slip system. The reduction of distortion can have the result that deformation continues in the primary system, which gets rid of a certain number of dislocations passing to the system. A characteristic feature of deformation in the third stage is the development of a cross-slip representing the main mechanism by which dislocations bypass the obstacles formed in the second stage.

Heat treatment processes for steels

There are three basic processes for the heat treatment of steels. These are (i) annealing, (ii) quenching, and (iii) tempering.

Annealing

Annealing process of steels has different methods namely (i) diffusion annealing, (ii) softening, (iii) phase recrystallization annealing or full annealing (normalization, high temperature or coarse grain annealing, and pearlitization), and (iv) stress relief annealing and recrystallization annealing.

The objective of diffusion annealing is to eliminate, as far as possible, in-homogeneities in the chemical composition, in particular liquation in-homogeneities, which occur during crystallization of steels. This annealing is usually carried out in the range of the gamma solid solution at a temperature of 1100 deg C to 1300 deg C. Diffusion annealing can be used primarily to smoothen out a difference in the content of alloying elements, the difference being due to the inter-crystal liquation. This shows up as smearing of dendrites with an increase in temperature and heating time. Differences in micro-hardness are removed simultaneously. The overall hardness of the steel decreases since liquation regions possessing high hardness is removed. Some average hardness is attained. The success of diffusion annealing largely depends on the steel purity and liquation. This type of annealing is generally used to improve properties of medium purity steels.

Softening is used to produce the structure of globular pearlite. This structure is very soft and readily lends itself to deformation during drawing and cold rolling etc. Steels with a low C content become too soft after this annealing treatment. The globular pearlite structure is favourable in steels with a C concentration of more than 0.5 %. Another goal of softening is to produce a uniform fine structure with finely dispersed C after quenching. The simplest method of softening consists in holding for many hours at a temperature slightly above Ac1 temperature. In this case, martensite which is left from the previous treatment is removed and the work hardening caused by cold working is eliminated. Cooling after softening can be done in air starting from 600 deg C. Refinement of the structure subjected to softening is achieved only above the point A1 temperature.

Phase recrystallization annealing consists of a twofold gamma to alpha transformation, which takes place during this annealing. It leads to the appearance of a fine grained uniform structure differing completely from the initial structure. Refinement of the grain during normalization results in the disappearance of the Widmanstätten and coarse grained cast structures, which have poor mechanical properties. Inhomogeneity of the structure in the work hardened state is removed. The closer the annealing temperature is to Ac3 temperature and the shorter the holding time at this temperature, the finer is the grain. Refinement of the grain structure is also facilitated if the heating rate to the annealing temperature and the cooling rate from this temperature are increased.

In the case of normalization, cooling is done in air. Here it is important to allow for different rates of cooling along the cross section of large sized products. The arising thermal stresses are removed by stress relief annealing or high temperature tempering. To obtain a fine grained structure, rapid cooling is done only over the transformation temperature interval. The normalization heating temperature is not to be much higher than the transformation point, or else the grain may be too coarse (overheating). An excessively long holding time also have the same result. The optimal heating temperature is determined by the C content.

Stress relief annealing and recrystallization annealing removes macroscopic stresses which are present in cold worked steels due to the dislocation pile ups and crystal lattice distortions. Normally these stresses are very high. Changes in properties which occur during the cold working can be rectified during subsequent heating. The greater is the degree of cold working, the lower is the heating temperature. Depending on the temperature and time of annealing, various structural changes take place in a cold worked material. The changes are divided into recovery and recrystallization processes.

Recovery is a totality of any spontaneous process of variation in the density and distribution of defects before the onset of recrystallization. If recovery proceeds without the formation and migration of sub-grain boundaries inside the recrystallized grains, it is called restoring. If sub-grain boundaries are formed and migrate inside the crystallites, recovery is referred to as polygonization.

Restoring does not include an incubation period. Properties start changing right at the beginning of annealing. Restoring is accompanied by a redistribution of point defects whose concentration decreases subsequently from excess concentration to the equilibrium concentration. Simultaneously, dislocations are redistributed and unlike-sign dislocations are annihilated.

The structure of steel starts changing drastically from a certain annealing temperature. New rather equilibrium grains are seen along with extended cold worked grains. They differ from the grains of the deformed matrix by having a more perfect internal structure. As distinct from the polygonized structure, recrystallized grains are separated from the matrix with large angle boundaries.

The formation and growth of grains with a more perfect structure which are surrounded by large angle boundaries at the expense of initially cold worked grains of the same phase is called primary recrystallization. Recrystallization begins with an incubation period. The recrystallization rate increases initially from zero to a maximum and then decreases due to an ever rising number of new grains in contact with one another.

Inclusions of insoluble impurities (carbides, nitrides) lower the tendency to growth of recrystallized grains. This is especially important in the case of ferritic steels, which are prone to grain growth. Another phase can precipitate during recrystallization in alloy steels which are subjected to a strong cold working.

Sometimes the intensive growth of individual crystals can be seen after a strong deformation and long holding (for several days) at temperatures close to the melting point. This phenomenon is called secondary or collective recrystallization.

Quenching

Quenching is the strengthening treatment. It consists of cooling from the temperature range of the solid solution at such a rate that transformation in the primary and bainite ranges are suppressed and martensite is formed. In this state, steel has the property of the highest hardness. There is a distinction between (i) normal quenching, which is used mainly for treatment of medium C and high C steels, and (ii) quenching after a thermo-chemical treatment (carburization, high temperature cyaniding etc.), which is used for low C steels.

In case of normal quenching for providing a required cooling rate during quenching, different cooling media and methods are used. Water, oil, or air can serve as the cooling medium. Many alloy steels, which have a high stability of austenite, are subjected to step quenching. With this method of quenching, the temperature drop is less than in the case of direct cooling to room temperature and consequently quenching stresses are less.

Some quantity of austenite is retained during quenching even in steels with a relatively small C content. For this reason it is impossible to impart the maximum hardness to a product. Since austenite is stable at room temperature and passes to martensite at lower temperatures, steels are being given a subzero treatment. Under this treatment quenching is continued and steels with a high content of retained austenite are immersed in liquid nitrogen (N2) or quenching mixtures whose temperature is below 0 deg C.

For surface quenching (if it is necessary to harden only the surface layer to a preset depth), special quenching heating schedules are used. The surface of the steel is fully heated, while the core is cold and remains unquenched on subsequent fast cooling. The selection of steel for surface quenching is governed by the sensitivity of the steel to quick heating and cooling. For this reason the C concentration is limited to 0.7 % to prevent cracking of the steel.

Among the quenching defects, the main defects are excessive holding and overheating. They show up as enlargement of martensite needles and coarse grain fracture. This leads to a high brittleness of quenched steel and the formation of cracks. Cracks are frequently formed at the boundaries of initial austenite grains. A low quenching temperature or too short a holding time at the given temperature causes incomplete quenching. In this case, the quenched steel is insufficiently hard.

Carburization is the thermo-chemical treatment which is associated with surface saturation of steel with C and N2. These elements quickly dissolve in Fe by the interstitial method and are capable of rapid diffusion to a considerable depth. Low C steels are subject to carburization. Carburization is usually carried out at 900 deg C to 950 deg C. Gas carburization is used mostly, under which steel is heated in the atmosphere generated from a gas which contains predominantly methane (CH4) or from liquid hydrocarbons (kerosene and gasoline etc.). Carburization is aimed at enrichment of the surface layer with carbon. The required strengthening of the surface layer is achieved by quenching, which is done after the carburization. The specific volume of the quenched carburized layer is higher than the specific volume of the core, and hence considerable compression stresses arise in the layer. This improves the fatigue strength of the steels.

Cyaniding consists of the saturation of the surface of steels with C and N2 in a cyanide-containing salt bath. The C-N2 ratio in the diffusion layer is controlled by changing the medium’s composition and the processing temperature. Advantages of cyaniding over carburization consist in a shorter processing time and improvement in the wear and corrosion resistance (owing to the presence of N2 in the surface layer).

Tempering

The main purpose of tempering is to provide a disperse structure at a preset degree of cooling. In the case of low C steels, quenching serves as tempering even if the steel is not subjected to high temperature tempering it has a high viscosity and a relatively high strength.

When some of the steels are quenched in oil, a structure is formed even during transformation in the bainite range which is more disperse than the one formed after cooling in air. But the most disperse distribution of carbides and the most favorable properties are obtained after martensite tempering. The structure dispersion has the highest effect on the YS. An improvement of the fracture stress and YS and an increase in the fracture stress -YS ratio can be taken as a measure of the tempering efficiency. The tempering efficiency depends on the cross-sectional area and on the content of C and alloying elements in the steels.

Although to achieve a thorough quenching the critical quenching rate has to be exceeded over the entire cross section, full tempering does not require this procedure. As an example, in quenched steel which has martensite in the surface zone and pearlite in the core, the hardness of the core can sometimes be higher than that of the surface zone after tempering. This is particularly the case during a short tempering when precipitation of carbides from martensite proceeds faster than the coagulation of pearlite plates.

Tempering of hypo-eutectoid steels which do not contain free ferrite gives a uniform improved structure. In the presence of ferrite precipitates, the fracture stress-YS ratio decreases and the impact strength is lesser than in the surface zone. Hence, in selecting the content of C and alloying elements and particular conditions of austenitization and cooling, the size of the steel product to be tempered is required to be considered. For tempering to give adequate properties, it is often enough to suppress the formation of ferrite during continuous cooling. Only when a very high fracture stress is required an abrupt cooling is used for tempering. In this case, susceptibility to full tempering can be improved by raising the quenching temperature and thus enlarging the austenitic grain size.